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OHMIC CONTACTS TO P-TYPE GALLIUM NITRIDE 



By 
BO LIU 



A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL 

OF THE UNIVERSITY OF FLORIDA IN PARTL\L FULFILLMENT 

OF THE REQUIREMENTS FOR THE DEGREE OF 

DOCTOR OF PHILOSOPHY 

UNIVERSITY OF FLORIDA 
2001 



'_ ACKNOWLEDGMENTS 

I would like to thank my supervisor, Dr. Paul H. HoUoway, for the opportunity 
and financial aid to study GaN, a promising semiconductor material. I am also grateful 
for his philosophy on the relation between research and knowledge, for his desire to 
understand students, and his devotion to improving my communication skills. I treasure 
the memory of his guidance and help with my research and personal affairs. 

■^ I thank my committee members: Drs. Rolf E. Hummel, Kelvin S. Jones, Fan Ren 
and Wolfgang Sigmund for their valuable contributions. The assistance from the staff of 
MAIC, Drs. Ren's group. Dr. Jones's group. Dr. Hummel's group and collaborators in 
the clean room in the Department of Electrical Engineering is highly appreciated. 

I thank all of the members in Dr. Holloway's group, both past and present. I 
appreciate their friendship and patience over the past few years. I enjoyed working with 

each of them. « <--. * 

I also thank all of the friends I made during my time in Gainesville, including 
those who favor me as their barber. I thank each of them for making my life in 
Gainesville so beautiful. 

I especially thank my family for their patience, support and encouragement 
through all my professional endeavors. 

This work was supported by EPRI/DARPA under agreement No.W-08069-07. 



n 



TABLE OF CONTENTS 

Eage 

ACKNOWLEDGMENTS • " 

vi 

LIST OF TABLES 

LIST OF FIGURES • ^" 

xi 

ABSTRACT 

CHAPTERS 

1. INTRODUCTION 

2. REVIEW OF LITERATURE ^ 

2.1 Growth of GaN and Defects ^ 

2.2 Mechanisms of Ohmic Contact ' 

2.3 Ohmic Contact to p-GaN: Present Research Status ^^ 

2.3.1 Fermi Level Pinning ^3 

2.3.2 Surface Preparation of GaN • 

2.3.2.1 Hydrogen ^5 

2.3.2.2 Carbon 26 

2.3.2.3 Oxygen .7 

2.3.2.4 Cleaning of Surface 

2.4 Interfacial Metallurgical Reactions 

2.4.1 Gallide-forming Metals 

2.4.2 Nitride-forming Metals 

2.4.3 Neutral Metals 

2.4.4 Thermal Stability • ^^ 

2.5 Metallization Schemes and Analysis 

2.5.1 Conventional Contact Schemes 

2.5.2 Non-conventional Contact Schemes 

3. EXPERIMENTAL PROCEDURES ^" 

46 

3.1 Introduction .^ 

3.2 Contact Preparation 

3.3 Characterization 



Ill 



4. EFFECTS OF H2O2 SOLUTION TREATMENT ONp-GaN 54 

54 



4.2 Modification of Electrical Conductivity 

4.2.1 Effects of H2O2 Concentration ^ 

4.2.2 Effects of Extended Immersion Time ^ 

4.2.3 Stability of the Increased Electrical Conductivity o^ 

4.3 Structural Characterization 

4.3.1 AES 

4.3.2 ESCA 

4.3.3 SIMS 

4.3.4 AFM 1: 

4.4. Application in Formation of Ohmic Contact to p-GaN ' j 

4.5 Discussion „- 

4.6 Summary 

5. "NOG" SCHEME FOR OHMIC CONTACT TO p-GaN 84 

84 

5.1 Introduction 

5.2 Principles of "NOG" Scheme °^ 

5.3 Comparison with Published Contact Results °° 

5.4 Experimental Studies 

5.4.1 Effects of Ti and Al as Nitride-Forming Metals ^^ 

5.4.2 Effects of Si and Mg as Nitride-Forming Metals • 9 / 

5.4.3 Neutral Metals j^^ 

5.5 Discussion 

5.6 Summary 

6. EFFECTS OF Ni CAP LAYER ON THIN Ni/Au CONTACTS TO p-GaN 112 

112 

6.1 Introduction 

6.2 Contact Electrical Properties J J^ 

6.2.1 Annealing Temperature 1 1^ 

6.2.2 Effects of Annealing Time J |^ 

6.2.3 Effects of O2 Flow Rate J J° 

6.3 Light Transmittance J^^ 

6.3.1 Effects of Annealing Temperature ^20 

6.3.2 Effects of Annealing Time J^l 

6.3.3 Effects of O2 Flow Rates J22 

6.4 Microstructure Characterization J^^ 

6.4.1 SEM J^^ 

6.4.2 AES Survey and Depth Profiling j^^ 

6.4.3 XPS Analysis J^^ 

6.5 Discussion 

6.6 Summary 



7 



CONCLUSIONS ^^^ 



IV 



162 

8. FUTURE WORK 

164 

LIST OF REFERENCES 

...175 
BIOGRAPHICAL SKETCH 



LIST OF TABLES 

Page 

LI GaN-based electric devices 

1 onH ralriilated values for the heat of formation AH^°' of 

32 



Table 

3 



2.1 Comparison of experimental and calculated values for the heat of formation AH^"^ of 
related gallides and nitrides 

2.2 Current metallization schemes of ohmic contact to p-GaN 

4.1 Atomic concentration of elements from AES surface survey analysis in MBE p-GaN ..63 

4.2 XPS results from a 1 :1, 300sec H2O2 cleaned p-GaN sample 66 

78 

4.3 Relation of hydrogen incorporation and processmg steps 

89 

5.1 Enthalpy and entropy of hydride formation 

92 

5.2 Electrical conductivity of selected nitndes 

5.3 Calculated driving forces of metal reactions to Ga and GaN 

5.4 Calculation of diffusion characteristic distance of selected metals in nickel HO 

6.1 Lattice constants (A) of components in oxidized Ni/Au contacts to p-GaN 147 

6.2 Values of surface tension of Ni and Au] 



4 i 



VI 



LIST OF FIGURES 



Page 
Figure 

2 
1.1 Relation of bandgaps and lattice constants 

o 
2.1 A nanopipe in GaN imaged by HRTEM 



2.2 Energy diagrams of a metal contact to a i 

12 



semiconductor ^" 

2.3 Change of the interfacial behavior factor S 

14 

2.4 Mechanisms of ohmic contact formation 

17 

2.5 Possible patterns used to measure contact resistance 

18 

2.6 Plot of measured resistance vs. contact separation 



19 

lucc jjiiao^^ v^vjMiiiu'n* 

2.8 Dopant locations inbandgap of GaN 



2.7 Pressure-temperature projection of a three phase equilibra 

21 



2.9 Barrier height vs. metal work functions on n-type GaN 

2.10 Calculated Ni-Ga-N diagram at 600°C 

2.11 Calculated (a)Zr-Ga-N and (b) La-Ga-N isothermal diagram at 298K 35 

2.12 Calculated phase diagram for Ti-Ga-N at 800°C 

2.13 Calculated (a) W-Ga-N and (b) Re-Ga-N diagram at 600°C 38 

49 

3.1 Configuration of contacts used in Hall measurement 

3.2 Schematic of light transmittance measurement 

4.1 Effects of H2O2 solution treatment on the I-V curves of 25A Ni/500A Ti/500A Au to 

MBEp-GaN 

4.2 Effects of H2O2 solution treatment on the I-V curves of 1 OOA Ni/500A Ti/500A Au to 

MBEp-GaN ^ 



vu 



58 

4.3 Hall measurement results of MBE p-GaN 

4.4 Effects of H2O2 solution treatment with extended time on the I-V curves of lOOA 

Ni/500A Ti/500A Au to MBE p-GaN 

4.5 Effects of H2O2 solution treatment with extended time 

4.6 Hall measurement results of H2O2 solution treated MBE n-GaN 61 

4.7 Effects of H2O2 treatment on MOCVD p-GaN 

4.8 Hall measurement on stability of H2O2 treated samples 

4.9 Comparison of XPS peaks from as-cleaned and 1:1, 300sec H2O2 treated GaN 65 

4.10 Negative SIMS depth profile for 5:1 H2O2 solution treated p-GaN 66 

68 

4.11 AFM images of GaN 

69 

4.12MicrostructureofMBE-GaN 

4.13 Effects of 5:1, 60sec H2O2 treatment on the I-V of 500A Ni/500A Au to p-GaN 70 

4.14 Relation among width of depletion region, carrier concentration and built-in ^^ 

potential in GaN 

4.15 Relation between pH values and H2O2 concentration 

4.16 Relation of carrier concentration decrease and nanopipe density in GaN 81 

5.1 Principle of "NOG" scheme 

5.2 I-V of Ni/Au, Ni/Ti/Au and Ni/Al/Au on p-GaN 

5.3 Effects of Ni thickness on the I-V of Ni/Ti/Au contact to p-GaN 97 

98 

5.4 Effects of thermal annealing on I-V data 

5.5 I-V curves of 500A Pt/500A Au, lOOA Pt/50A Si/500A Pt/500A Au andlOOA Pt/50A 

Mg/500A Pt/500A Au contact on p-GaN, as-deposited 

5.6 I-V curves of 500A Pt/500A Au, lOOA Pt/50A Si/500A Pt/500A Au and lOOA 

Pt/50A Mg/500A Pt/500A Au contact on p-GaN, 600°C for 1mm annealmg 99 

5.7 I-V curves of 500APt/500AAu, 100APt/50ASi/500APt/500A Auand lOOA 

Pt/50A Mg/500A Pt/500A Au contact on p-GaN, 800°C for 1mm annealmg 100 

5.8 AES depth profile of Pt/Au contact on MOCVD p-GaN 1^2 



Vlll 



99 



1 0*^ 

5.9 AES depth profile of Pt/Si/Pt/Au contact on MOCVD p-GaN 

5.10 AES depth profile of PtMg/Pt/Au contact on MOCVD p-GaN 10^ 

5.11 1.Vof500ANi/500A Au, lOOANi/lOOOATi/lOOOA Au and lOOA Ag/IOOOA ^^^ 

Ti/IOOOA Au on p-GaN 

5.12 Comparison of In metal and lOOA Ni/500A Ti/500A Au as ohmic ^'''''^^'^^^^^^q^ 
measurement 

6.1 Effects of anneal temperature on I-V of 50/50 contact 

6.2 Effects of amiealing temperature on the specific contact resistance of the 50/50, ^ ^^ 

50/50/50 and 50/100/50 contacts 

6.3 Effects of annealing time on the specific contact resistance of 50/50, 50/50/50 and ^ ^^ 

50/100/50 schemes annealed at 600°C 



6.4 Effects of anneal time on resistance of contact pads at 600°C 

121 



120 

6.5 Comparison of light transmittence at X = 450 nm 

1 99 

6.6 Effects of annealing time on Hght transmittance at X = 450 nm • 

6.7 Effects of O2 flow rates on Ught transmittance at 500°C. 

6.8 Microstructure of the 50/50 contact after annealing. 

6.9 Microstructure of the 50/50/50 contact after annealing 125 

6 10 SEMbackscattering electron image of same sample as in Figure 6.9-(d) but at a 

higher magnification showing the Au film is still continuous 1^0 

6.11 Microstructure of the 50/100/50 contact after annealing 127 

128 

6.12 EDS analysis of the light region in Figure 6.16-c 

128 

6.13 EDS analysis of the dark region in Figure 6.16-c 

129 

6.14 EDS analysis of spherical particle in Figure 6.16-d 

130 

6.15 Annealed Au film on GaN and sapphire 

131 

6.16 Annealed (600°C, lOmin) Ni film on GaN 

132 

6.17 AES surface spectra fi-om 50/50 contacts 

133 

6.18 AES surface spectra fi-om 50/50/50 contacts 



IX 



134 

6.19 AES surface spectra from 50/100/50 contacts 

135 

6.20 AES depth profile of 50/50 contacts 

6.21 AES depth profile of 50/50/50 contact 

137 

6.22 AES depth profile from 50/100/50 contacts 

6.23 XPS spectra of Ni2p from 50/50 contact after annealing at 600°C 139 

6.24 XPS spectra of Ni2p from 50/50/50 contact after annealing at 600°C MO 

6.25 XPS spectra of 01s from 50/50 contact after annealing at 600°C 1^1 

6.26 XPS spectra of 01s from 50/50/50 contact after annealing at 600°C 142 

6.27 XPS spectra of Ga2p from 50/50 contact after annealing at 600°C 143 

6.28 XPS spectra of Ga2p from 50/50/50 contact after annealing at 600°C 144 

6.29 Energy diagram of oxidized thin Ni/Au contact to p-GaN 1 

6.30. Schematic diagram of interface equilibrium between three phases 148 

6.31 AE/p (= rp/R) for p = 42 and 10/5 calculated for 50/50 and 50/50/50 152 

6.32 Schematic diagram of contact microstructure 






Abstract of Dissertation Presented to the Graduate School 

of the University of Florida in Partial Fulfillment of the 

Requirements for the Degree of Doctor of Philosophy 

OHMIC CONTACTS TO P-TYPE GALLIUM NITRIDE 

By 

Bo Liu 
May 2001 
"'^"r DeTa "meT^^^^^ Science and The effects of H.O. treatment, multi-layer 
metallization, and Ni cap-layer on Ni/Au have been studied for ohmic contacts to p-GaN. 
First, surface H2O2 treatments are found to increase the hole concentration by up to 100% in 
p-GaN grown by molecular beam epitaxy (MBE), while causing no change in n-GaN or p- 
GaN grown by metalorganic chemical vapor deposition (MOCVD). Treatment of 20 min 
increased, while treatment >60 min decreased the hole concentration in MBE p-GaN. With 
this treatment, the current in Ni/Au contacts increased. The increased hole concentration was 
attributed to reduction of nitrogen vacancies or H-Mg complexes in GaN. The decrease of 
carrier concentration was attributed to recompensation of shallow acceptors by oxygen 

serving as hole traps. 

Second, general principles were defined for selecting metals for ohmic contacts to 
GaN, a scheme called "NOG" for Nitride-forming metal Over Gallide-forming metal. In 
"NOG" a gallide-forming metal dissociates GaN and a nitride-forming metal increases the 
nitrogen thermodynamic activity at the interface. Literature data were compared to these 
ideas, and experimental data on Ni/Ti/Au, Ni/Al/Au, Pt/Si/Pt/Au, PtMg/Pt/Au were 
collected and compared with data from Ni/Au and Pt/Au contacts. Higher currents were 



XI 



found for schemes based on the "NOG" principles. However, the contact resistivity was still 
high and thermal stability became a limiting factor. 

. ■ Last, ohmic contacts to p-GaN were obtained after oxidizing Ni/Au and Ni/Au/Ni 

contacts. Both Ni/Au/Ni and Ni/Au were shown to have resistivities of ~1 Q-cm . 
Transparent NiO was obtained and thin Au film formed pores which led to optical 
transparencies atX = 450nm of >85%. The porosity in Au was demonstrated to result from 
interface and grain boundary energies. Addition of the Ni cap-layer was shown to increase 
the thermal stability of thin Ni/Au ohmic contacts and increase the light transmittance to 
93%, while keeping contact resistivities of low 10"* Q-cm . 






v'f,-*-' •■? 



xu 



CHAPTER 1 
INTRODUCTION 

As members of the III-V nitrides family, InN, GaN, AIN and their alloys are all 

wide band gap materials, and can crystallize in both hep wurtzite (a) and cubic 

zincblende (P) crystal structures. As shown in Figure 1.1, the wurtzite InN, GaN and AIN 

have direct bandgaps of 1.9 eV, 3.4 eV and 6.2 eV, respectively, at room temperature 

[Mor97]. In the cubic form, GaN and InN have direct bandgaps, and AIN has an indirect 

bandgap. The GaN alloyed with hiN and AIN can form a continuous (AlGaln)N alloy 

system spanning a continuous range of direct bandgap energies throughout the visible to 

near UV region of the electromagnetic spectrum. This makes the nitride system very 

attractive for optoelectronic device applications, such as light emitting diodes (LEDs), 

laser diodes (LDs) and optical detectors [Mor96]. Various GaN-based electric devices 

also have been demonstrated for applications in high power/high temperature electronic 

devices [Table 1.1] because of their intrinsic properties of wide bandgap and high 

breakdown fields [Cho94, Kha94, She99b Gas98]. 

In particular, GaN with its band gap of 3.4 eV plays a central role in the alloy 

system of (AlGaln)N. Other advantageous properties of GaN include high mechanical 

and thermal stability, large piezoelectric constant, good thermal properties [Bin97] and 

the possibihty of passivation by forming thin layers of GaiOs or AI2O3 with bandgaps of 

4.3eV and 9.2eV respectively. 



1 

CQ 




a-InN 



P-AIN 



1 1 1 1 1 1 1 ■ ' 




300 K 



^ 1 1 ■ 1 1 ■ I ■ ■ I »■ I ■ ■ ■ ' 1 1 ■ ■ ■ I ' ' ' ' 1 ' ' ' 



3 3.2 3.4 3.6 3.8 4.2 4.4 4.6 4.8 S.O 

Lattice Constant (A) 

Figure 1.1 Relation of bandgaps and lattice constants of hexagonal (a) and cubic (P) InN, 
GaN, AIN and their alloys [Mor97] 



Because of its wide bandgap and chemical inertness, GaN and related materials 
are challenging materials researchers. Examples of some challenges include but not 
Umited to growth of heteroepitaxial films [Amb98], the low level of p-type doping 
because of deep acceptor levels [Mol93] and Mg-H complex [Got95], difficulties in 
achieving low-resistance ohmic contact to p-GaN [Liu98], slow wet and dry etch rate 
[Ade93, Pea94], and lack of good passivation films [Uza95], etc. This work focuses on 
formation and improvement of ohmic contacts to p-type GaN. 



Table 1.1 GaN-based electric devices [Shu99] 



Status 



Device 

Schottky barrier diode Demonstrated for a variety of metals 

p-n junction Demonstrated on both regular and 

lateral epitaxial overgrowth 
material 

Demonstrated 



GaN MESFET 



GaN MISFET 



AlGaN/GaN HFET 



AlGaN/GaN and 
GaN/SiC HBTs 



Possible Applications 



Demonstrated 



Demonstrated on sapphire and SiC 
substrates with record power levels 

Demonstrated 



GaN-based piezoelectric Demonstrated for GaN and 
and piezoelectric AlGaN/GaN 



sensors 



GaN pyroelectric sensor Demonstrated for primary and 

secondary pyroelectric effect 



Switch, FET building block 

Photodetector, switch, BJT 
building block 



High-temperature digital 
circuits 

High-temperature digital 
circuits, non-volatile 
memories 

High-temperature, high power 
microwave, high temperature 
digital circuits 

High-temperature, high power 
microwave, high power 
switches 

Pressure sensors, especially for 
high temperature 
applications 

Temperature sensors, especially 
for high temperature 
applications 



Progress has been made in the past several years in developing reliable ohmic 
contact to p-GaN [Tre96, Tre97, Hol97, Liu98, Ho99a, Ho99b, Koi99 and PalOO], both 
for lower specific contact resistance and better thermal stability. Nevertheless, the goals 
are far from being achieved. It has been proven difficult to obtain sufficiently low contact 
resistance (<10-^ Q-cm^) to p-GaN. In GaN based LEDs, Ni/Au are commonly used as 
ohmic contact on p-type GaN top layer. The low doping levels of the p-GaN usually 
result in non-ohmic contact, and thereby degrade the device performance. Also, due to 
the high resistivity of the p-GaN layer, the current from the top electrode can not be 
spread effectively through the entire device chip, leading to current crowding. Large 



Joule heating caused by high operating voltages has been reported to limit the lifetime of 
continuous wave (CW) laser diodes [Nak97]. For metal semiconductor field effect 
transistors (MESFETs), an ohmic contact resistivity of mid -10'^ C^-cm may be 
acceptable for a channel length of 1 ^m [Mur90]. For p-GaN, the best stable ohmic 
contact resistivity is only lO' ~ lO"' Q-cml Ohmic contacts to n-GaN have been 
developed with resistivities of 10'^ Q-cm^ [Ren97]. 

The high specific contact resistances to p-type GaN can be attributed to several 
factors, including 1) absence of a metal with a sufficiently high work function (The 
bandgap of GaN is 3.4 eV and the electron affinity is 4.1eV, so a work function of 7.5 eV 
is needed for a good ohmic contact. Metal work functions are typically < 5.5 eV); 2) 
difficulty in achieving high hole concentrations in p-GaN because of the deep ionization 
level of the acceptors (Mg is -170 meV, others are deeper [Str91]); and 3) the tendency 
for preferential loss of nitrogen firom the GaN surface during processing, which probably 
result in surface conversion to n-type conductivity. To decrease the contact resistance to 
P-GaN, high p-type electrical conductivity or lower barrier height at the contacts interface 
would be helpful. . 

Specific contact resistivifies as low as 10"^ Q-cm^ have been reported for thin 
oxidized Ni/Au films [Ho99b]. This contact scheme also showed high (>60%) light 
transmittance, which is desired in optoelectronics. The reported contact thickness (only 
100 A) was comparable to the GaN surface roughness. Lower contact resistivity and 
higher light transmittance are obviously desirable. 

In this dissertation, H2O2 solutions were demonstrated to increase the carrier 
concentration of MBE grown p-GaN. A "NOG" scheme (Nitride-forming Over Gallide- 



forming metals) offers guidance in development of contacts to both n- and p- type GaN 
epilayers. Low resistance and high transparency contacts to p-type GaN are also studied. 
It was demonstrated that adding a thin Ni surface cap-layer, to the thin Ni/Au layered 
contacts improved light transmission and enhanced contact stability. 
Thus the scope of this dissertation is as follows. 

Chapter 2 presents a review of the physical mechanisms of ohmic contact 
formation and discusses current literature on ohmic contact to p-GaN. 

Chapter 3 summarizes the experimental procedures used in this study, including 
contact processing parameters and characterization techniques. 

Chapter 4 summarizes the use of H2O2 solutions to increase the free hole 
concentration. Possible explanations for increased hole concentration are discussed. 

A methodology to choose materials to test for better GaN ohmic contact 
formation is proposed and discussed in Chapter 5. This procedure ("NOG") is based on 
the properties of different metal groups. Various contact schemes to p-GaN are tested and 

discussed. 

In Chapter 6, the microstructural evolution of transparent Ni/Au ohmic contacts is 
reported. The consequence of annealing these layers in N2 or O2 ambients are described. 
Addition of a Ni cap-layer to form GaN/Ni/Aumi contacts is shown to increase rather 
than decrease the optical transparency. 

Finally, Chapter 7 is a summary, and chapter 8 discusses future work to improve 

the understanding of p-GaN ohmic contacts. 



' CHAPTER 2 

REVIEW OF LITERATURE 

Many GaN-based electronic and photonic devices have been developed, and all 
these devices need ohmic contacts to external current and voltage sources. Because of 
high conductivity, metals are normally used for this purpose. Typically, interfaces 
between metals and semiconductors have been found to exhibit rectifying properties 
[Rho88, Bri93]. Although these rectifying properties are useful in some devices, many 
applications require efficient transport of current across the contact interface. Ohmic 
contacts are generally required for such applications. Creating low-resistance ohmic 
contacts to GaN, especially for p-type material, has proved to be one of the major 
challenges faced in the development of GaN semiconductor technology [Pea97a]. 

2.1 Growth nf GaN and Defects 
Because sufficiently large (> 1 cm in diameter) single GaN crystals are generally 
unavailable for use as substrate for homoepitaxial growth, heteroepitaxial growth of GaN 
has been used in practice and the choice of substrate is critical to the structural, electrical 
and morphological properties of the obtained epilayers [Li96, Kob97, Ham98, Kun96, 
Kur95, Geo96, Sun96, Pop97]. Possible substrate materials typically need to have good 
matches in thermal expansion coefficients and lattice constants with GaN, and should be 
resistant to the growth chemistries at high growth temperatures (over 1000°C in certam 
cases). Sapphire [Kai98, Pop97] and SiC [Pop97, Abe97] are the most popular substrates 
currently used due to their adequate thermal and chemical stability at high growth 

6 



temperatures, excellent structural and surface morphology and availability in large 
quantities. Other substrates include Si [Vis95], MgAl204 [Kur95], LiGa02 [Kun96], 
NdGa03[Kur95], quartz glass [Iwa97] and ZnO[Pop97,Dav97]. ^ 

Metalorganic chemical vapor deposition (MOCVD) [Ama86, Kat94, Kel96, 
MorSl and Nak94], molecular beam epitaxy (MBE) [Mou93, Van97] and their 
derivatives are the techniques most extensively used to grow GaN epilayers. Mg and Si 
are widely used as p - and n - type dopants. Because of a large activation energy (270 
meV) and passivation of acceptors with hydrogen, in which the activation energy is 
dependent on dielectric constant (GaN, 9.5) and effective mass of the material (GaN, me 
= 0.2 mo, mhh = 0.75 mo) [Den97], the ionization of Mg acceptors is less than 1% at room 
temperature. Typical hole concentrations are ~10" cm"^ for p-GaN (although Mg 
concentration can be as high as 10^° cm"^). For GaN grown with the MOCVD method, 
the as-grown materials typically is unintentionally n-type, which is widely believed to be 
due to intrinsic nitrogen vacancies. Intentional n-type doping can be easily 
accompoHshed using silicon as the donor. 

hi the MOCVD method, the Ga source materials generally used are GaCb 
(produced by passing HCl vapor over molten gallium), trimethylgallium (TMGa), 
triethylgaUium et al, and the nitrogen source is mostly of ammonia [Dav88, Wal97]. 
Biscyclopentadienyl (CpzMg) is used as a source of Mg and methyl silane (Si (CHsSiHa) 
as a source of Si [Sun97, Her97]. The MOCVD method has been the leading technique 
for production of Ill-nitrides optoelectronic and microelectronic devices. Characteristics 
of this method include high purity chemical sources, easy composition and uniformity 
control, high growth rates and abrupt junctions. Since partial pressure of nitrogen in the 



CVD reactor is always less than the equilibrium partial pressure of nitrogen over GaN at 
the high substrate growth temperatures, the samples may contain high concentrations of 
intrinsic n-type carriers, commonly believed to be nitrogen vacancies [Pan97]. 

In the MBE process, the fluxes of Ga, Mg and Si are generated by heating high 
purity elements in the Knudsen or effusion cells [Dav88]. The substrate temperature for 
MBE growth is typically operated at relatively low temperature at 898K or even lower to 
663K [Got81]. Nitrogen is typically supplied as an atomic species using electron 
cyclotron resonance (ECR) or radio frequency (rf) plasmas [Pop98, Abe97, Mou93]. 


















.■9 m #9#*###^*< 






Figure 2.1 A nanopipe in GaN imaged by HRTEM [Kan99] 



The GaN-based Ill-nitride heterostructures are found typically to contain characteristic 
one-dimensional (edge, mixed and screw dislocations) and two-dimensional (stacking 
faults and domain boundaries) extended defects [JaiOO]. Although the dislocation density 
is high, the dislocations are usually clustered in local regions of the epilayer, so large 
volumes of the epilayer are defect free [Pon97]. Nanopipes are also found in GaN with 
diameters ranged from ~ 5nm to ~ 0.5um and densities as high as 10* cm'^ as identified 
with high-resolution transmission electron microscopy (HRTEM). These nanopipes were 
reported to be parallel to the c-axis of GaN unit cell [Ven99, Kan99], as shown in Figure 
2.1. The main constituents inside these nanopipes were Ga, C and O [Kan99]. 

■^ A model was constructed to explain the effects of dislocations on minority 
carriers in GaN epilayers [Jai98], and was found to be consistent with the experimental 
results. Dislocations and nanopipes are postulated to help explain changes in the carrier 
densities of MBE grown GaN after room temperature treatment with H2O2 solutions in 
this work. 

2.2 Mechanisms of Ohmic Contact 
A contact is said to be ohmic when the ratio of the potential drop F across a 
contact versus the current / flowing through the same contact is linear with a constant, 
low contact resistance {Re). Ohmic contacts are characterized by a parameter called 
specific contact resistance (or resistivity), pc, which is expressed as [She92] 



10 




(a) 






Ef 

Evb 



*. 



.^L...t 



w . 



Semiconductor 



<t>M 



-,- -Etn 



Contact 



(b) 



Figure 2.2 Energy diagrams of a metal contact to and semiconductor, 

explaining formation of ohmic contact to p-GaN (a) Before charge 
equilibrium; (b) After charge equilibrium. 



In general, when a semiconductor of a given electrochemical potential is brought 
into contact with any phase with a different electrochemical potential, charge will flow 
automatically across the semiconductor/contact junction (Figure 2.2). For an ideal 
semiconductor/metal contact (Schottky limit), all the voltage drops across the 
semiconductor. The contact barrier height for p-type semiconductor, ^, p, is calculated 



using the equation [Kum93]: 

n,p = 



(2.2) 



where E/, „ is the Fermi level of the isolated metal (before contact) and x is the electron 
affinity of the isolated semiconductor. In the literature, the work fiinction of a metal, (p (in 



11 

electron volts), is often used to estimate the barrier height limiting charge transfer at 
semiconductor/ metal junctions. The work fiinction is used instead of Ef_ „ because the 
Fermi level is more difficult to determine experimentally, whereas ^is readily accessible 
using photoemission or other data. Ideally, a metal with a lower work function than that 
of an n-type semiconductor or a higher work function than that of a p-type semiconductor 
can be used to form ohmic contacts to a semiconductor. Unfortunately, it has been shown 
experimentally that most semiconductor/metal contacts do not obey the predictions of 

this ideal Schottky limit [Rho88, Kum93]. hi a simple model, the Schottky barrier height, 

^, can be expressed as follows [SzeSl]: 

q-A-^(S-%^+</>o) (2-3) 

where %„ is the metal electronegativity (Note: not to be confiised with Xs, the electron 

affinity of the semiconductor), and (po represents the contribution of surface states fi-om 

the semiconductors. The interface index 

"■■ ' ■ ' "' dz„ 

is a function of the electronegativity difference zl;ir between the cation and anion 
components of a compound semiconductor, as shown in Figure 2.3* [Ren98]. The reason 
for the name of "5" factor is just fi-om the shape of the graph. Note the sharp transition 
around Ax=l. This S factor should be equal to unity for an ideal semiconductor/metal 
junction in the Schottky limit. When 5 = 1, the measured barrier height equals the initial 
contact potential difference obtained from the Schottky limit. The unity of 5 value means 



• The electronegativity difference of GaN was labled mistakenly at 1.8 eV in the literature, instead of the 
true value of 1.23eV as shown in Figure 2.3. 



12 



1 


ZnS 




AIN SiOj 


ZnO SrTiOj 


« 1 - 

1^ 




T 


__t.^_ 


1 1 1. _ - «■* 

Al,03 KTa03 


o 

1 0.8 - 

& 

1 0.6- 


f GaS 
1 

■ / ■ 

CdS' 

^icaSe 






Bcpected 
Value 


c 

i 0.4 - 


SiC / 
Gap /"Se 








- 0.2 - 


• / 

Ge GaAs'' Actual 


4 


n-type 




L.llSb |.'^p Values 




p-type 




n -5' "^ 


^ 






4 1 1 

0.5 1 




1.5 


2 




Bectronegativity Difference 



2J 



Figure 2.3 Change of the interfacial behavior factor S with different 
semiconductors [Ren98]. 



absolutely no pinning of the Fermi level. When 5-0, the system is in the regime of 
strong "Fermi level pinning." This terminology indicates that the Fermi level position at 
the surface of the semiconductor, measured relative to the vacuum level, does not vary 
when either the work function or the Fermi level of the contacting phase is changed. For 
covalent semiconductors with Az<\,S is small, and (^ is typically affected by a high 
density of surface states from dangling bonds, so that it depends weakly on the metal 
work function. On the other hand, for ionic compound semiconductors, where Ax>l, the 
index S approaches unity, and ^ depends strongly on the metal work function. GaN has 
an electronegativity difference of 1.23 eV(Ga: l.SleV, N: 3.04eV) [HsuOO], which 
suggests that the Schottky barrier height should be a function of the metal work function. 
The origin of this non-ideal interfacial behavior (Fermi level pinning) is a topic of 
intense controversy. Bardeen originally proposed [Bar47] that surface states at the 



13 

semiconductor/metal interface are the source of this pinning. Chemical reactions, such as 

stoichiometry changes [Spi86] and alloy formation [FreSl], have also been proposed to 

explain Fermi level pinning. 

For either an ohmic or rectifying contact, current is transported across the 

interface by mechanisms shown in Figure 2.4. Current transport can be principally 

ascribed to the following three mechanisms [SzeSl]: 

1 . Thermionic emission (TE) is dominant in low and moderately doped 

semiconductors with Ne(h) ^ lO'"' cm"l At low to moderate carrier densities, 
the wide depletion region prevents tunneling through the barrier. When the 
barrier height is small, the electrons can be thermally excited over the top of 

the barrier (thermionic emission. Figure 2.4a). On the other hand, for a high 

barrier, the vast majority of electrons are unable to overcome this barrier, 

'• '- ' . resulting in non-ohmic (rectifying) contacts. 

\ -. 2. Thermionic-field emission (TFE) is applicable for intermediate doping 

f 
■ ' ' densities, -lO'"' cm'^ < Ne(h) < ~10'^ cm"^ Both thermionic and tunnel emission 

are important (Figure 2.4b). 

3. Field emission (FE) is effective in heavily doped semiconductors, Ne(h) > ~10 
cm'''. In this case the depletion region is narrow, and electron or hole tunnel 
easily from metal to semiconductor or vice versa (Figure 2.4c). 

A very useful parameter, KT/Eqo, can be used [Yan71] to calculate pc for each of 
these three mechanisms, where Eoo is the tunneling parameter equal to: 



£„=f^j^ (2-5) 



14 





u 


^e.. 


> 






(a) 


V 



«t>b 


'./. 


"^^^^^^^ eeee' 


> 






(b) 


V 



" e"*e" 



(c) 






p f p '^ "- 



Figure 2.4 Mechanisms of ohmic contact formation, (a) thermionic 
emission; (b) thermionic field emission; (c) field emission. 



15 

where q is electronic charge, h is the Planck's constant, Ne(h) is electron (hole) 
concentration, tis dielectric constant of the semiconductor, and wi* is electron (hole) 

effective mass. 

For kT/Eoo « 1 (heavy doping concentrations), the specific contact resistance is 

given by 

■^00 

In this case, the field emission mechanism (FE) dominates current transport. The pc 
depends strongly on doping concentration. At high doping concentration, the depletion 
width of the Schottky junction is decreased, resulting in high tunneling transmission 
coefficients. Hence, even a metal with a low metal work function can still form an ohmic 
contact. With this method, as in GaAs, ohmic contacts can be formed on a semiconductor 
with a pinned Fermi level [Hol97]. 

For ^r/£oo~ 1 (intermediate doping concentrations), a mixture of both thermionic 
and tunneling (TFE) transport is observed, and the specific contact resistance becomes 






p, oc exp 



^■<i>b 



£, 



£^.coth(-^) 



(2.7) 

The specific contact resistance in this case depends on both temperature and transmission 
coefficient for turmeling. 

For kT/Eoo » 1 (moderate doping concentrations), the TE mechanism dominates 
the current conduction and the specific contact resistance is 

p,ocexp(^) (2-8) 



16 



The specific contact resistance is clearly dependent on temperature. At higher 
temperatures, the thermionic emission current increases and results in a smaller A- 

The term ohmic contact in practice does not necessarily require a linear current 
voltage characteristic [Rid75]. A metal-semiconductor contact is associated with a space- 
charge region in which the current-voltage behavior eventually becomes nonlinear, as 
bias increases. Ideally the contact resistance of the space-charge layer would be 
negligible relative to the bulk or spreading resistance of the semiconductor contacted by 
the metal, but this is rarely achieved in practice. A contact is usually acceptable when the 
voltage drop is very small compared to the drop across the active region of the device, 
even though the current-voltage behavior of the contact is not strictly linear. This is 
certainly true for current contacts to p-GaN devices, due to the unavailabihty of low 
resistance ohmic contact to p-GaN (see later in Table 2.2). 

In theory, the contact resistance can be defined completely if physical and 
operating parameters are known [Rid75]. The physical parameters are mainly of contact 
area and thickness, while operating parameters are predominantly temperature and bias. 
In practice, the contact resistance can be affected seriously by a number of other factors, 
such as interfacial layers (oxide formation or contamination), surface damage, minority 
carrier injection, and energetically deep impurity levels or traps. 

The most widely used method for determining the specific contact resistance is 
transfer length method, also often called the transmission line model (TLM) [Ber72, 



17 



(•) 




(b) 



Figure 2.5 Possible patterns used to measure contact resistance: (a) 
linear array; (b) circular contacts [Wil84]. The grey region is 
contact metal and the white region presents areas where the 
metal has been etched or lifted off leaving bare semiconductor 
or an etched mesa. 



Wil84]. An array of metal contacts (Figure 2.5) is fabricated with different spacings 
between the metal areas. The resistance is measured as a function of the gap spacing. 
Extrapolation of the resistance to zero gap spacing gives a value equal to twice the 
contact resistance Re (Figure 2.6). The x-intercept is equal to twice the transfer length, L,, 



18 



C/3 

Pi 



T3 /* 



<U 



/• 



2Li 



dL W 



2R. 



Contact spacing 



Figure 2.6 Plot of measured resistance vs. contact separation [Wil84] 



where L, = (pM "\ pc is the specific contact resistance and /?, is sheet resistance of the 
semiconductor epilayer. The transfer length is defined as the distance from the edge to 
where the current in the semiconductor falls to 1/e {e being the base of natural logarithm) 
of its original value, pc can be calculated using the equation: 



Pc = 



R 



= Rs'L'. 



(2.9) 



where L, should be much smaller than the gap spacing between contacts. However, as the 
spacing become very small, irregular edges may cause a problem. The test pattern should 
be isolated so that current flow occurs only in the space between pads (no leakage current 
path), therefore a mesa structure may be needed [Wil84]. 



19 



500 



T.oC 




g 10 

iOOOO/T(K) 



Figure 2.7 Pressure-temperature projection of the three phase equilibra: GaN + 
liquid + gas. The area inside each envelope represents the two-phase 
equihbrium of GaN + gas and corresponds to GaN stability [Dav99]. 



For the pattern in Figure 2.5-(b) with a circular TLM (CTLM), device isolation 
can be omitted because there is no leakage path for the current flow. The total resistance, 
R,, calculation is complex, but when r/I, » 1, it can be calculated using [Ho99b]: 
R. 



R.= 



1-n 



ln(-) + A-4 + -) 
r R r 



(2.10) 



where Rs represents the sheet resistance of p-GaN, R and r denote the radii of the outer 
and inner circular contacts, respectively, and L, is the transfer length. The total resistance 
is measured for different spacings and plotted as a function of ln(R/r). The least square 
curve-fitting method can be used to obtain a straight linear plot ofR, vs ln(R/r). The slope 
gives Rs, and the intercept at ln(R/r) = is Rs-L/r-K, leading to U so the specific contact 
resistance, pc, can be expressed as 



L. = m 



Pc\l/2 



(2.11) 



20 

In practice, linear I-V curves cannot always be obtained. To quantify the degree 
of linearity conveniently, an arbitrary figure of merit [Piq98] is defined as follows: 

LM = 



dr'~-' (2.12) 



dl 



v=0 



where the derivatives are the resistance at 5V and OV, respectively, and Z,M ranges fi-om 
zero to unity, approaching unity for samples with nearly linear I-V curves. 

A good ohmic contact should have a low specific contact resistance, high 
stability, smooth surface morphology and good edge definition. Because contact 
resistance usually constitutes a small portion of the total measured resistance, caution 
must be taken to avoid errors. 

Metallic contact layers usually are prepared by vacuum deposition (electron beam 
or thermal evaporation) followed by heat treatment. The most commonly used method of 
heat treatment for the metal-semiconductor system is furnace annealing in H2, N2, or N2 + 
H2 forming gas to protect the contact metals fi-om oxidation. For p-GaN, the annealing 
environment should not contain H2, due to the compensation of Mg acceptors with H 
element [Sug98]. 

2.3 Ohmic Contact to p-GaN: Present Research Status 
Initially, as grown GaN films were unintentionally n-type with a high carrier 
concentration. This is widely believed to be due to nitrogen vacancies being intrinsic 
donors [Neu94]. Vacancies result from the large vapor difference between the Ga and N 
components as shown in Figure 2.7 [Dav99] The nitrogen vacancy lies just below the 



21 



Conduction Band ,, .. 

Ga site '^ ^"^ 



-30meV '^ 

Native Defect Level 



750 meV Li 

700 meV Be 

550 meV Cd 

410 meV Hg 

340 meV Zn ^25 meV Si 

250 meV ' Mg 



Valence band 
Figure 2.8 Dopant locations in bandgap of GaN [Den97] 

conduction band (~ 30 meV), which makes it an efficient donor (Figure 2.8) [Str91, 

Ben97]. 

In contrast to n-type, a critical breakthrough was achieved when p-type GaN was 
reported using Mg dopants followed by low electron beam irradiation (LEEBI) [Aka89]. 
It has been shown that interstitial hydrogen is incorporated into GaN to form an H-Mg 
acceptor complex, thus passivating the Mg acceptors. The H-Mg bonds can be broken 
with LEEBI or high temperature annealing in an inert ambient [Pea96]. Acceptor doping 
using Be has also been predicted [Ber97]. To understand ohmic contact formation, the 
subjects of pinned Fermi level, surface cleaning, interfacial metallurgical reactions and 
data on multilayer metallization schemes are discussed in this section. 



22 



2.3.1 Fermi Level Pinning 

Semiconductors can be classified into two groups based on the dependence of 

Schottky barrier height on metal work function. The ionic materials have a direct 

dependence of barrier height on metal work function whereas covalent materials have a 

weak dependence, as discussed in section 2.2 and shown in Figure 2.3. 

From the ionic nature of GaN (an electron negativity difference of 1.23eV 

between Ga and N atoms), the Schottky barrier height of GaN is expected to depend 

directly on the metal work function. The Schottky barrier heights change with the metal 

worlc function, but the changes are much less than expected for both doping types of 

GaN. Mori et al [Mor96a] found that the change in Schottky heights is smaller than the 

difference between the metal work function and the work function of p-GaN for Pt, Ni, 

Au and Ti metals. The Schottky barrier heights on n-type GaN are also found to depend 

weakly on the metal work function [Guo95]. hi both cases, deposition was made by 

electron beam evaporation and the contacts were not subjected to heat treatment after 

deposition. The metal work function was not the dominant factor affecting the Schottky 

characteristics of Pt and Pd on n-GaN. Rennie et al studied the electrical properties of 

various metal (Ti, Al, Sn, Cu, Zn, Mo, Ni and Pd) contacts to n-type and p-type GaN in 

the hope of determining the relationship between metal work functions and barrier 

heights [Ren98]. Contrary to the expected trend, the Fermi level was calculated to be 

strongly pinned, with the effect being greater in the p-type material (Figure 2.9). The S 

factor is shown to be only 0.01 for p-type and 0.21 for n-type GaN (in Figure 2.3), in 

contrast to the expected 5=1 for a completely unpinned Fermi level (Schottky limit). 



23 



1.2 



a> 



0.4 



4.2 



4.8 5 5.2 

Work Function, eV 









♦ 




• 






X 






A 






R 




5.4 


5.6 5.8 



Figure 2.9 Barrier height vs. metal work functions on n-type GaN from various 
researchers [CaoOO] 



- Another basic consideration is the sum of the p- and n-type Schottky barrier 
heights for the same metal, which is found to be significantly less than the bandgap of 
GaN. The measured values of Schottky barrier heights on p-GaN are likely to be affected 
by the presence of interface states, damage and contamination. More research is 
necessary to understand the interfacial conditions and properties of metal/GaN contacts. 

2.3.2 Surface Preparation of GaN 

The preparation of the GaN surface before metallization influences the electrical 
characteristics of metal/GaN contacts [Kin96]. Cleaning the surface in solvents and 
common acids or bases is effective in removing a significant fraction of the surface 
oxides and other contaminants, but these procedures can not produce atomically clean 



24 

surfaces. Surface roughness, which can vary depending on the techniques and conditions 
used to grow GaN epilayers, can influence the uniformity of contacts. One important 
issue in the processing of ohmic contact to GaN is to understand the semiconductor 
surface and to find appropriate chemical treatments to clean it before contact metal 
deposition. Ideally, the metal/semiconductor contact should be oxide- and defect-free, 
atomically smooth, uniform, and thermally stable. 

The most common impurities on the surface after preparation of ohmic contact to 
GaN include C, H and O fi-om the ambient air, alcohol, methanol, acetone, water, 
photoresist residues and sample handling, etc. [Abe96]. Their effects are discussed 
below. 

2.3.2.1 Hydrogen 

A small amount of H2 in the carrier gas can passivate the electrical activities of 

Mg and C acceptors in p-GaN during cooling-down from MOCVD growth [Sug98]. This 

reduces the p-type doping levels unless an annealing step is performed. Annealing at 450 

~ 500°C can restore the fi-ee hole concentration but hydrogen atoms do not physically 

leave the films until a higher (>800°C) temperature is reached [Pea96]. Low resistivity p- 

type GaN could be obtained by H2-free MBE growth without any post-treatment [Sug98]. 

Hydrogen is predicted to act as a donor (H"") in p-type GaN, and as an acceptor 
(H') in n-type material [Neu96]. Implantation of ^H^ creates high resistivity materials 
fi-om both n- and p- GaN [Pea98]. 

For ohmic contacts to p-GaN, the Mg-H complex may be dissociated at any 
period during processing. In principle, dissociation of the Mg-H complexes before or 



25 

after metallization should have no impact on final resistance. But for Pt/Au, Pt, Pd/Pt/Au 
and Ni/Au contacts to p-GaN [Kin97], lower a values were obtained in the absence of 
premetallization annealing. The post-metallization anneal simultaneously activates the p- 
dopant and anneals the contact and/or contact interface. Annealing in N2/O2 mixture 
environment decreased the contact specific resistance for ohmic contact to p-GaN by a 
factor of three [Suz99]. This is attributed to O2 reacting with H in GaN to form H2O and 
uncompensated Mg acceptors. Some researchers disagree with this postulated 
mechanism. 

2.3.2.2 Carbon 

Although carbon has been shown to produce p-type GaN, the hole concentration 

obtained has been limited to lO'^ cm'^ even though the carbon concentration was 10 

cm'^ or higher [Pan76]. It has been found in other III-V materials that the maximum hole 

concentration fi-om carbon doping is related to the difference in bond strength between 

the group III- and group V-carbon sites, hi the case of InP, the carbon actually sits on the 

group III site and acts as a donor resulting in an n-type material. This is also expected to 

occur in InN and high In concentration alloys of InxGai.xN and hixAli-xN grown by 

MOMBE. Carbon has been shown to be a strongly n-type doping element for x > 0.15 in 

hixGai.xN and X > 0.3 in KAli-xN. As the In concentration is reduced, the tendency 

increases for carbon to act as an acceptor rather than a donor. It has also been proposed 

[Abe96] that C displays amphoteric doping behaviors in the nitride with acceptor 

formation under some conditions (MOMBE grown GaN) and donor formation in other 

cases (implantation in GaN and growth of In containing alloys). 



26 



2.3.2.3 Oxygen 

From photoluminescence data [Sei83], it was concluded that oxygen is neither a 

shallow nor a very deep donor. It has a moderately deep level (about 78 meV below the 

conduction band edge at 4.2 K) which could form an impurity band near the conduction 

band edge. 

When O is implanted into GaN and annealed at 1 100°C [Abe96], it creates n-type 
doping with an ionization level of -29 meV. Seifert et al. [Sei83] proposed that 
substitutional incorporation of oxygen onto nitrogen sites could be the origin of intrinsic 
free carriers in the growth of highly conductive n-type GaN. Because of the similarity of 
their atomic radii, both substitutional oxygen and nitrogen vacancies could cause donor 
defects. 

Using an (Al, Ga, hi) bubbler to purify the NH3 during the growth of GaN, a 
dramatic reduction of carrier concentration in n-GaN is found due to the removal of 
H2O/O2 from the NH3 [Chu92]. Seifert et al. [Sei83] used Mg3N2 to purify the NH3, and 
also observed the reduction of as-grown carrier concentrations. Further verification of the 
contribution of H2O/O2 to the increase of the carrier concentration comes from 
experiments where water is introduced intentionally during growth. The carrier 
concentrations of the water-injected samples are always above 10 cm" . 

Chemisorption of oxygen on atomically clean and ordered GaN(OOOl) surfaces 
showed that saturation occurs at coverage of 0.4 monolayers [Ber96]. Low energy 
electron diffraction (LEED) indicates an ordered adsorbate layer, and x-ray photoelectron 
spectroscopy (XPS) peak of O Is suggests a single chemically distinct adsorption site. 



27 

The oxygen also is reported to react with GaN to form monocHnic P-Ga203 [Wol97] at 
900°C when GaN (both film and powder) is exposed to dry air. An interfacial reaction 
mechanism is identified as the rate limiting step for oxidation, with an apparent activation 
energy of -72 kcal/mole. The oxidation resulted in roughening of the oxide/GaN 
interface and oxide surface. 

2.3.2.4 Cleaning of Surface 

As stated earlier, the metal/semiconductor interface should be inert, oxide- and 
defect- free, atomically smooth and covered by epitaxial metal. For GaAs and other III-V 
semiconductors, a thin oxide layer (~3 to 20 A) grows rapidly on the surface when 
exposed to air, necessitating in-situ cleaning of the GaAs surface under ultra-high 
vacuum (UHV) for the epitaxial growth of metal films. 

For GaN, in-situ deposifion of Ga metal followed by thermal desorption under 
ultra-high vacuum is found to yield atomically clean surfaces using Auger electron 
spectroscopy (AES) [Kha93]. This has been used in the study of Ni [Ber93] and Al 
[Ber96b] films on GaN. hi-situ nitrogen ion sputtering and annealing can also produce 
atomically clean GaN surfaces [Ber96a, Hun93]. 

Ex-situ cleaning is usually used in practical metal deposition. Although solvent 
cleaning and wet etching with common acids or bases cannot produce atomically clean 
surfaces, they are effective in removing a significant fi-action of the surface oxides and 
other contamination. This results in relatively intimate metal/GaN interfaces. The effects 
of aqua regia (HNO3: HCl = 1 :3), HCl: H2O, HF: H2O and NH4OH and NaOH for 
cleaning the surface of GaN are also investigated [Kin96]. A HCl based solution is found 



28 

to be more effective in removing oxides and leaving less oxygen residue, but HF is more 
effective in removing carbon and/or hydrocarbon contamination. The HCl and HF based 
solution should be equally effective in removing the total contamination. 

The importance of surface preparation is exemplified by cleaning in boiling aqua 
regia [Kim98] for 10 min and then depositing a layered structure of 200 A Pd/5000 A Au 
in a vacuum of 10''' Torr on p-GaN (Nh= 2.98 x lO'^ cm'^). A good contact resistance, pc 
= 4.3 X 10"^ Q-cm^ was obtained on as-deposited samples. With no surface treatment, 
samples deposited with the same metallization exhibited a high resistance of 2.1x10' Q- 
cm^ The lower contact resistivity is attributed to removal of the surface oxide from the p- 
GaN surface. Similarly, with a short time interval between the GaN film growth and 
metal film deposition, a good ohmic contact is also possible as obtained in Ref. [Jan99] 
even if the samples are only ultrasonically degreased with trichloroethylene, acetone and 
methanol for 5 min each step. A contact resistivity of -3x1 0'^ Q-cm^ is found for as 
deposited p-GaN/Pt/Ni/Au samples.and a value of 5.1x10""* Q-cm^ is reported after 
annealing at 350°C for 1 min in an inert ambient. 

2.4 Interfacial Metallurgical Reactions 
Interfacial reactions are critical to the formation of ohmic contacts to 
semiconductors, whether they have a large or a small bandgap. Interfacial reactions can 
lead to disruption of interfacial contamination layers consisting of native oxides, 
hydroxides, and hydrocarbon/organic residue layers due to reaction or adsorption 
[Hol97]. Although some work has been reported on contact schemes to p-GaN, little 
information is available about the metallurgical reactions on the metal/p-GaN system. 



2^ 

Calculations about the transition metal-Ga-N systems have been performed on the 
metallurgical reactions of metals with GaN in the absence of experimental studies 
[Moh96]. According to the enthalpy of reactions to form gallides and nitrides, all 
transition metals can be classified into three groups, the late, early and middle transition 
metals. They correspond to the gallide-forming, nitride-forming and neutral metals 
discussed for the "NOG" contact scheme in Chapter 4: 

2.4.1 Gallide-forming.Metals 

These are mainly group VIII metals. The group VIII metals are characterized by 
the absence of intermediate phases in the metal-N binary systems with either positive or 
small negative enthalpies of formation (Table 2.1). For Ru, Rh, Pd, Ir and Pt, no metal 
nitrides have been reported, although these systems have not been investigated 
thoroughly. The nickel nitrides are believed to be metastable under 1 atmosphere or lower 
N2 pressures at and above room temperature, and osmium nitrides are not considered 
because no thermodynamic data are available. In contrast to nitrides, group VIII metals 
form many metal gallides, like NiGa, Ni2Ga3, etc. 

Two types of tie line configurations are predicted for metal-Ga-N phase diagrams 
for these gallide-forming (late transition) metals. Figure 2.10 shows the Ni-Ga-N tie-line 
configurations with different N2 pressures as an example. Changing the pressure of N2 
represented at the top comer of the isothermal phase diagram can alter the tie-line 
configuration. In an annealing environment, N2 is also predicted to have effects on how 
far the metal/GaN reaction can be driven in these systems. 



30 



N2(latm) 



r2 , 



Ga(l) 



NiGa 
(a) 



N2(10-'atm) 




Ni 



Ga(l) 




N2(2xl0-^atm) 




Ga(i) 



(c) 



Ni 



Figure 2.10 Calculated Ni-Ga-N diagram at 600°C. The nitrogen comer of the 

diagram represents (a) N2 at 1 atm; (b) N2 at 10"^ atm and (c) N2 at 2 x 
10-^atm [Moh96] 



31 

Higher reaction temperatures increase driving forces for reactions between these 
metals and GaN. At 600°C, a tie-triangle between NiGa, GaN and N2 gas at 1 atmosphere 
is observed (Figure 2.10-a). Assuming that a Ni contact is much thinner than the 
underlying GaN layer, Ni/GaN is favored to react under 1 atmosphere of N2 at 
thermodynamic equilibrium. The entire contact should be converted to NiGa/GaN, and 
nitrogen gas is released during the reaction. In the calculation, it was also shown that a tie 
line exists between Ni2Ga3 and GaN in the phase diagram. However, the equilibrium 
partial pressure of N2 over a Ni2Ga3-GaN contact would be less than 1 atmosphere. Thus, 
there would be a driving force for a Ni2Ga3 contact to react with N2 at 1 atmosphere to 
form NiGa and GaN. Of course, the reaction would be too slow to be observed, just as in 
the case of GaN, which does not grow appreciably on liquid Ga at 600°C under 1 
atmosphere of N2, even though the reaction is very thermodynamically favorable. 

The prediction of reactions between Ni and GaN is supported by experimental 
data. The growth of thin Ni films on GaN [Ber93] is examined and a pronounced reaction 
occurrs upon annealing above 600°C in vacuum. Ga4Ni3 is identified by x-ray diffraction 
(XRD) along with Ni in the as-deposited film [Guo96]. Although no reaction is found for 
Ni on GaN in as-deposited state, after annealing at various temperatures between 400 and 
900°C, a trend of increasing Ga in the reacted films is observed with increasing 
temperature [Ven97]. New phases consistent with NisGa and NiGa are found. 

The tie-line configuration predicted for these metal-Ga-N diagrams is strongly 
affected by the stability of the metal gallides, the temperature of the isothermal section 
and the N2 pressure of interest. Thermodynamic data shows that Pd and Pt gallides are 



32 



Table 2.1 Comparison of experimental and calculated values for the heat of formation 
AH*^"' of related gallides and nitrides. Published experimental values for the 
entropy of formation, AS*^"', have been added when available. The units for AH °' 
and AS*""' are kJ/mol and J/(K-mol) respectively [Boe88] 



System 


Compound 


AH^o^exp 


Ml'\,^c 


AS'°^ 


Sc-Ga 


ScGa 




-68 





Sc-N 


ScN 


- 157 (T unknown) 


-184 





Ti-Ga 


TiGa 


. _ . 


-51 





Ti-N 


TIN 


-169 (298 K) 


-146 


-48 






,. - 173 (298 K) 


_ 




V-Ga 


V^Gas 


-11(763- 953 K) 


-18 


- 3.8 (763- 953 K) 




VeGas 


-16(763- 953 K) 


-28 


- 5.1 (763- 953 K) 




VjGa 


- 19 (763 - 953 K) 


-20 


-13.8 (763- 953 K) 


V-N 


VN 


-109 (298 K) 


-76 


- 45 (298 K) 




V2N 


- 90 (298 K) 


-74 


-33 


Cr-Ga 


CrGa4 


-7.6 (850 K) 


-8 


+ 2.9 (850 K) 




CfsGae 


-4.7 (850 K) 


-15 


+ 4.8 (850 K) 




CrjGa . 


-3.3 (850 K) 


-12 


+ 2.7 (850 K) 


Cr-N 


CrN 


- 53 (298- 1800 K) 


-22 


- 35 (298- 1800 K) 




CrzN 


-31 (298- 1800 K) 


-35 


- 17 (298- 1800 K) 


Mn-Ga 


MnGa 


" .' 


-34 







Mn3Ga2 




-34 





Mn-N 


Mn5N2 


- 34 (298 K) 


-44 







MatN 


- 26 (298 K) 


-31 





Fe-Ga 


FeGaj 


- 20 (298 K) 


-10 







FcgGai 1 


- 33 (298 K) 


-16 







FevGae 


- 35 (298 K) 


-18 





Fe-N 


FezN 


-1.3 (298 K) 


-11 







Fe4N 


-16 (298 K) 


-7 









-2.2 (298 K) 











- 2.4 (298 - 860 K) 








Co-Ga 


CoGaa 


-45 (298 K) 


-18 









-29 (HOOK) 









CoGa 


-41 (298 K) 


-31 


— 



Table 2.1 -Continued 



33 

















-32 (298 K) 


__^ 









-36 (1173 K) 




-5.4(1 173 K) 






-38 (1173 K) 





— 


Co-N 


C03N 


+ 2 (298 K) 


+ 1 




Ni-Ga 


NijGaT 


- 34 (300 K) 


-37 


— 




NizGaj 


-45 (300 K) 


-33 






NiGa 


-38 (300 K) 
- 47 (298 K) 


-37 


— 






-43 (1023 K) 


_ 








. -43 (1223 K) 





-2.9 (1223 K) 




Ni3Ga2 


-36 (300 K) 


-37 








. -45 (298 K) 


_^ 




Ni-N 


NijN 


+ 0.2 (292 K) 


+ 6 




Ru-Ga 


RuGa 




-34 




Ru-N 


RuN 




+ 49 





Rh-Ga 


RhGa 




-53 




Rh-N 


RhN 




+ 46 


_ 




RhsN 




+ 11 





Pd-Ga 


PdjGa 


-59 (1000 K) 


-51 







PdggGaia 


-27 (1000 K) 


-24 





Pd-N 


PdN 




+4« 


_ 




PdsN 




+ 12 




Os-Ga 


OsGa 


- 28 (298 K) 






Os-N 


OS5N 




+ 12 





Pt-Ga 


PtjGa 


-44(1000K) 


^M 







Pt94Ga6 


- 6.7 (lOOOK) 


-11 




Pt-N 


PtN 




+ 62 


_ 




PtsN 


\ 


-!« 


— 



34 

particularly stable among the group VIII metal gallides. The reaction of these metals with 
GaN is thermodynamically favorable. They form metal gallides and release N2 gas, even 
under 1 atmosphere of N2 at room temperature, although these reactions may be 
extremely slow at this temperature. 

2.4.2 Nitride-forming Metals 

Be, Mg, Ca, Sr, Ba, Al and transition metals like Sc, Ti, V, Y, Zr, Nb, La, Hf and 
Ta are nitride-forming metals. Si can also be classified into this group because of the 
large enthalpy of the Si-N reaction. The majority of the non-transitional metal-nitrides 
have low electrical conductivity, hence only the transition metals are considered [Table 
2.1]. In contrast to the late transition metals, these early transition metals form metal 
nitrides of considerable thermodynamic stability. These metal-N binary systems all 
contain a refractory metal mono-nitride (MN phase). Some of them actually exist over a 
wide range of compositions and some systems contain nitride phases besides mono- 
nitride. 

In all of the systems for which no ternary phases have been reported (Sc, Y, Zr, 
La and Hf metals), the common feature in the calculated phase diagrams is a tie line 
between the MN phase and GaN. Such a tie line would also be expected for the (Nb, Ta, 
V, and Ti)-Ga-N systems, as long as tie lines to any ternary phases do not alter this 
situation. At room temperature, the predicted MN-GaN tie lines are stable against 
competing tie lines by at least 40kJ/g-atom, due to the enhanced thermodynamic stabiHty 
of the MN phases. Even at 600°C, these tie lines still represent a negative enthalpy of 
reaction of more than 15kJ/g-atom. Figures 2. 11 -a and b. show the calculated Zr-Ga-N 



35 



N2(latni) 



N2(latm) 




Ga(l) 



Zr 




Ga(l) 



LaG% 



La 



Figure 2.1 1 Calculated (a)Zr-Ga-N and (b) La-Ga-N isothermal diagram at 298K 
[Moh96] 



and La-Ga-N diagrams. To simplify the diagrams, the ranges of homogeneity of the 
binary phases are neglected. Taking ZrN as an example, there is not enough data to 
predict the exact compositional range in equilibrium with GaN. In contrast to the results 
of gallide-forming metals, the temperature and N2 pressure represented in the nitrogen 
comer of the diagram have less dramatic effects on the calculated phase diagrams, at least 
over the range of temperature and pressure normally encountered in contact processing. 

The phase diagram for the Ti-Ga-N system at 800°C is shown in Figure 2.12 
[Moh97]. According to this diagram, if a N2 pressure of 1 atm is maintained continuously 
over the contact, the Ti/GaN contact would be thermodynamically favorable to extract 
nitrogen from the annealing environment, ultimately resulting in a TiN/GaN contact. This 
result is expected because only TiN and GaN are simultaneously in equilibrium with N2 
gas at 1 atm. The contact would therefore come to equilibrium with GaN without any net 
consumption of GaN through an interfacial reaction. However, there may be competition 



36 



N2(latm) 



Ga(l) 



H H H H H H H H H jj 



Figure 2.12 Calculated phase diagram for Ti-Ga-N at 800°C [Moh97] 

between nitrogen incorporation into the film and metallurgical reaction at the non- 
equilibrium Ti/GaN interface. Because the Ti contact surface is usually covered by a 
protective metal layer (like Au), the reaction of metals v^ith the annealing environment is 
usually hampered. A Ti contact would react with GaN to ultimately form TiNx and leave 
liquid Ga remaining on the GaN. This situation is more favored when contacts are 
annealed under a lower N2 partial pressure. 

Reactions of metals with GaN to form nitrides plus Ga is expected to decrease the 
hole concentration at the metal/p-GaN interfacial region due to a subsequent increase of 
N vacancies in the interfacial region. However, before the formation of stable nitride, 
these metals could allow more nitrogen atoms to diffuse into the semiconductor/metal 
interface. This can decrease the nitrogen vacancy concentration and is expected to 



37 

increase the hole concentration. So, the contact could be expected to exhibit low 
resistance at short time, but degrade to higher resistance at long time after the formation 
of stable nitrides. Literature data are discussed relative to this concept in Chapter 5. 
Specifically, Ta/Ti metallization forms low resistance contact to p-GaN, but degrades 
after a few days [Suz99]. This is also believed to be the reason that Ti, Al and Ti/Al 
forms good ohmic contact to n-GaN [Les96, Cor98, Ruv96]. Interfacial reactions should 
also contribute to ohmic contact of W and WSix to n-hiGaN and n-hiN [Var97a, Var96a, 
Var97b]. 

2.4.3 Neutral Metals 

Neutral metals include Cr, Mn, Mo, Tc, W and Re. Cr and Mn can form 
compounds to both Ga and N (see appendix A). The Mo-Ga, W-Ga and Re-Ga binary 
phase diagrams are characterized by an absence of intermediate phases under 
atmospheric pressure and negligible miscibility between liquid Ga and the metals. Figure 
2.13 shows the W-Ga-N and Re-Ga-N diagrams at 600°C. Both W and Re are expected 
to be in thermodynamic equilibrium with GaN at room temperature and 600°C. Although 
W2N, WN and ReNo.43 are reported to form, none of these nitrides are expected to be 
stable at 600°C under 1 atmosphere or lower N2 pressures. Mo is predicted to be more 
like a gallide-forming metal because a tie-triangle is predicted at 600°C between a metal 
gallide, GaN and N2 gas at 1 atmosphere in the Mo-Ga-N diagram. 

The Cr is found to form a contact with linear I-V curves to p-GaN after annealing 
at 900°C for 15 sec [Tre97]. This is speculated to result fi-om the reaction of Cr with both 
Ga and N. This dual reaction may improve the carrier concentration in the metal/p-GaN 



38 



N2(latni) 



N2(latm) 





Ga(l) 



W Ga(l) 



Re 



(a) 0>) 

Figure 2.13 Calculated (a) W-Ga-N and (b) Re-Ga-N diagram at 600°C [Moh96] 

contact region under some condition, but this improvement may vary in performance 
with the change of temperature and time. 



2.4.4 Thermal Stability 

An increase of operating or processing temperature would allow Mg-H complexes 
to dissociate (to reactivate the p-type dopants), thus the carrier transport via thermionic 
emission could increase as temperature increased. The kinetics of interfacial reaction may 
be also increased. All these effects would improve the ohmic contact quality. 

The temperature is shown experimentally to have a strong influence on the 
contact properties. Detailed studies of the electrical properties of the Pt/Au contacts to p- 
GaN revealed that the I-V linearity improves significantly with measurement at higher 
temperatures [Kin97]. At room temperature, a slightly rectifying I-V curve is observed, 
while at 200°C and above the I-V curve is linear. When temperature is increased from 
25°C to 350°C, the specific contact resistance is found to decrease by nearly one order of 



39 



magnitude. A minimum pc of 4.2x10-^ Q-cm^ is obtained for a Pt/Au contact at 350°C. 
The behavior of Pt, Pd, Ni is also studied on n-GaN as a function of annealing 
temperature [Dux97]. The Pt film began to form submicron spheres and islands after 
annealing above 600°C, and the Pd and Ni films changed their morphology to islands for 
Pd/GaN and Ni/GaN interfaces, respectively. Delamination occurred at Pd/GaN 
interfaces upon annealing above 700°C, but no delamination occurred at Ni/GaN 
interfaces because of a SiOa cap layer on the Ni film. No structural changes are observed 
below these temperatures using XRD and RBS analysis. 

The delamilation of metal contacts on GaN was explained with the concept of 
thermal stress generated by differences in thermal expansion coefficients between metals 
and GaN. The surface energy and linear coefficient of expansion (6x10" K' ) of GaN are 
lower than most metals. Since the thermal expansion coefficient of Pd (11 .7x10 k ) is 
almost twice that of GaN, the change between room temperature and the annealing 
temperature generates an extension stress causing delamination. It may be necessary to 
use metals with low surface energies and low thermal expansion coefficients to avoid this 
delamination problem. Stable, non-reactive contacts that can withstand high temperatures 

are desired. 

It is believed that the current is governed by thermionic emission in current p- 
GaN samples with carrier concentration level of lO'^ cm'l The improvement in I-V 
linearity at high temperature is attributed to the increase in thermal energy of carriers 
which enables them to activate over the barrier. On the other hand, because the large 
ionization energy for Mg, -250 meV as shown in Figure 2.8, the percentage of activated 
Mg acceptors is low. The increase of operating energy would increase the activation 



40 

percentage of Mg acceptors, and improve the hole concentration. Increased carrier 
density could lead to field emission or thermionic field emission as the dominant 
mechanism of charge transport. As temperature increases, the I-V linearity improves and 
the contact resistance decreases significantly. At temperature above 200°C, the I-V curve 
exhibits ideal ohmic behavior and the resistance is constant w^ith current. 

2.5 Metallization Schemes and Analysis 
Of significant interest in the improvement of GaN devices is a p-type ohmic 
contact of low resistance. Common reported values are around 10" ohm-cm but a lower 
limit has been reported of 3x10"^ Q-cm^ for a doping level of 7xl0'^cm"^ [Suz99]. 
Recently, contact schemes with better performance have been demonstrated. These 
include oxidized thin Ni/Au [Ho99b, Koi99] and polarization-charge-based contacts 
[LiOO]. In this section, the current published contact schemes are compared and the 
mechanisms analyzed. Based on these contact mechanisms, the contact schemes are 
classified into conventional or non-conventional methods. Conventional methods mean 
formation of contacts via interfacial reactions to improve doping levels in the contact 
region in the GaN epilayer. Non-conventional methods used innovative designs to 
increase the carrier concentrations other than by interfacial metallurgical reactions. 

2.5.1 Conventional Contact Schemes 

These contacts fall into the scheme of the "NOG" concept developed below in 
Chapter 5. They will be analyzed in the discussion section in that chapter. In Table 2.2, 
the contact schemes, processing condition and contact resistivities are summarized. 







41 






Table 2.2 Current metallization schemes of ohmic contact to p-GaN 




Metalization (*) 
. Au250 


Pc, Q-cm^ 


Nd(cm"') 


Comment 


Refer 


53 


1x10'* 


800°C,10mininN2 


[Sim96] 


AulOO 


0.026 


MglxlO'° 


AD, 1mA bias 


[Mor96] 


Au32/Mp32/Aul70 


214 


1x10" 


Linear AD, more resistive 
w/750°C,15secinN2 


[Smi96b] 


Co80/Aul00 


8.1x10-^ 


1.6x10'^ 


AD 


[Fun99] 


Cr20/Au300 


1.2x10"' 


1.4x10^° 


500°C, 1 min in N2 


[Yoo96] 


Cr20/Au300 


2x10' 


2.7x10" 


700°^ ImininNi 


[Yoo96] 


Cr20/Au300 


1x10-^ 


1.4x10^° 


AD 


[Yoo96] 


Cr50/Aul00 


<4.3xl0' 


9.8x10"* 


900°C, 15secinN2 


[Yoo96] 


In200/Aul00 


1.5x10"' 


1.6x10'^ 


500°C, 1.5min 


[Fun99] 


Inl5/Zn20/In55/Aul00 


1.7x10"' 


3.1x10'^ 


RT~500°C, 1.5 min 


[Fun99] 


Mel 20/ AulOO 


Bad 


1.6x10'^ 


Up to 850°C 


[Fun99] 


Mn50/AulOO 


Bad 


1.6x10" 


Up to 850°C 


[Fun99] 


Nil 00 


0.015 


Mg 1x10-° 


AD, 1mA bias 


[Mor96] 


Ni80/Aul00 


1.8x10"' 


1.6x10" 


500°C, 1.5min 


[Fun99] 


Ni80/Aul00 


3.1x10"^ 


3.1x10" 


500°C, 1.5min 


[Fun99] 


Ni20/Au500 


3.4x10"' 


0.1-1x10" 


500°C, 30 sec 


[Kim97a] 


NiiO/Au40 


-10-^ 


8.4x10'^ 


500°C,10mininvac 


[Ish97] 


Nil0/Au5 


<10"' 


2*10" 


400-500°C, lOmin in O2 


[Ho99a] 


Ni5/Au5 


4x10"* 


2*10" 


500°C, lOmininair 


[Ho99b] 


Ni80/Aul00 


3.8x10"' 


2.2x10'^ 


500°C, 1.5min 


[Fun99] 


TSIi?/Au? 


1.2x10"^ 


5.5x10" 


lmAbias,700°C, lOmin 


[Kin97] 


Nil5/Crl5/Au500 


-8.3x10" 


' 0.1-lx-lO 


' 500°C,30secinN2 


[Kim97a] 


Nil5/Crl5/Au500 


8.3x10"^ 


0.1-lx-lO 


'^ AD 


[Kim97a] 


Ni25/Mg5/Ni25/Si240 


9.6x10"" 


3x10'^ 


400°C, lOmin PR/30 min 
RTAinN2 


[Kam98] 


Ni25/Mgll/Ni25/Si240 


6.6x10"' 


3x10" 


400°C, lOmin PR/30 min 
RTA in N2 


[Kam98] 


Ni25/Mg8/Ni25/Si240 


2.2x10"' 


3x10" 


400°C,10min PR/30 min 
RTA in N2 


[Kam98] 













Table 2.2- continued 



42 



Ni25/Mg8/Ni25/Si240 


2.5x10"' 


3x10'' 


450°C, 30 min in N2 


[Kam98] 


Ni25/Mg8/Ni25/Si240 


4.5x10"' 


3x10" 


500°C, 30 min in N2 


[Kam98] 


Ni8/Znl32/Aul00 


2.5x10"' 


3.1x10" 


400°C, 1.5 min 


[Fun99] 


Ni45/Zn-Au46 


3.6x10"' 


4.4x10" 


600°C, 2 min in N2 


[You98] 


Ni2.5/Znl0/Inl0 


2.3x10"' 


2.2x10" 


600°C, 1.5min 


[Fun99] 


/Ni47.5/Aul00 










Pd20/Au500 


9.1x10"' 


9x10" 


500°C,30secinN2 


[Kim97b] 


Pd20/Au500 


9.1x10' 


9x10'* 


500°C,30secinN2 


[Kim97b] 


Pd20/Au500 


2.9x10"^ 


2.98x10" 


AD, no clean, 3x10"' Torr [Kim98] 


Pd20/Au500 


4.3x10"^ 


2.98x10" 


AD, 


[Kim98] 


Pd20/Au20/Pd20/Au500 


-9x10"' 


9x10" 


AD 


[Kim97b] 


Pd20/Au20/Pd20/Au500 


-9x10' 


9x10" 


500°C, 30 sec in N2 


[Kim97b] 


Pd20/Au20/Pt20/Au500 


-9x10"' 


9x10" 


AD 


[Kim97b] 


Pd20/Au20/Pt20/Au500 


Mid 10"^ 


9x10'* 


500°C, 30 sec in N2 


[Kim97b] 


Pd?/Pt?/Au? 


1.0x10"^ 


5-6x10'* 


lmAbias,700°C, 10 min 


[Kin97] 


Pd?/Pt?/Au? 


1.7x10"^ 


5-6x10'* 


1mA bias, AD 


[Kin97] 


Pt? 


3.4x10"^ 


5.5x10'* 


1mA bias, AD 


[Kin97] 


PtlOO 


1.3x10"^ 


MglxlO^° AD, 1mA bias 


[Mor96] 


Pt?/Au? 


5.7x10"' 


5.5x10'* 


750°C, lOmin, cap/GaN 


[Kin97] 


Pt20/Au300 


-1.8x10"' 


1.4x10^° 


AD 


[Yoo96] 


Pt20/Au300 


2x10"' 


2.7x10" 


700°C, Imin in N2 


[Yoo96] 


Pt/Au 


4.2x10"^ 


5-6x10'* 


Measured at 350°C 


[Kin97] 


Pt20/Ni30/Au80 


-3x10' 


3x10" 


AD 


[Jan99] 


Pt20/Ni30/Au80 


5.1x10"" 


3x10" 


350°C, Imin in N2 


[Jan99] 


Ta50 


>10"' 


7x10" 


800°C, 20minA^ac 


[Suz99] 


Ta60/Ti40 


3x10"' 


7x10" 


800°C, 20minA^ac 


[Suz99] 


Ti50 


>10"' 


7x10" 


800°C, 20minA'ac 


[Suz99] 


Til 00 


3.5x10"^ 


MglxlO 


=° AD, 1mA bias 


[Mor96] 


Ti20/Pt80/Au300 


4x10"" 


1.4x10-° 


AD 


[Yoo96] 


Ti20/Pt80/Au300 


2.5x10' 


2.7x10'* 


700°C, ImininNz 


[Yoo96] 


2n50/Aul00 


4.9x10"^ 
unit: nm 


1.6x10" 


750°C, 1.5min 


[Fun99] 


Metallization thickness 






?: Not reported; 










AD: As Deposited. 


" ■ ' • -' 









43 



2.5.2 Non-conventional Contact Schemes 

Recently the use of thin Ni/Au to p-GaN for both low resistance and high- 
transparency ohmic contact has been extensively studied [She99b, Koi99, Ho99a, Ho99b, 
Che99, Mae99]. Transparent NiO is a p-type semiconductor with a wide bandgap, which 
varies from 3.6 to 4.0eV [Sat93]. Lower specific contact resistance values (<1.0xlO Q 
cm^ for lOOA Ni/50AAu and later of 4x10"^ Q-cm^ for 50A Ni/50A Au) are obtained 
[Ho99a, Ho99b]. Some Au-based contacts (Ni/Au, Co/Au, Cu/Au, Pd/Au and Pt/Au) 
annealed in an O2 partial pressure are also found to reduce the contact resistance, pc, and 
the sheet resistivity of p-GaN epilayers (a) [Koi99]. 

While significant progress has been made on ohmic contacts to p-GaN by 
oxidizing the thin Ni/Au, different mechanisms have been suggested to explain the 
results. Murakami's group [Koi99] ascribed the low contact resistance to formation of an 
intermediate semiconductor layer (ISL) with high hole concentration caused by removal 
of hydrogen atoms which were bonded with either Mg or N atoms in the p-GaN epilayer. 
Ho et al [Ho99a, Ho99b, and Che99] explained the results as being due to the formation 
of a p-NiO layer directly on the p-GaN layer surface. This NiO layer is believed to act as 
a low barrier ISL and small islands of Au bridged the current path via electron tunneling. 
This microstructure resulted in the reduction of the Schottky barrier height (^) at the p- 
GaN/metal interface. However, Murakami's group [MaeOO] further studied a variety of 
Ni/Au based contacts, such as 500A NiO/500A Au, 500A NiO (Li)/500A Au, 50A 
Ni/500A NiO (Li)/50A Au, 50A Ni/200A LizO/SOOA NiO/500A Au and 50A Ni/200A 
Li2O/50A Ni/500A NiO/500A Au (much thicker than the 50A Ni/50A Au in Ho's work), 



44 

using a method of sputter deposition and annealing in an O2 ambient for 5min in the 
temperature range of 300 to 500°C. They found that all contact schemes with NiO mid- 
layer did not reduce the pc to values lower than the conventional Ni/Au contact annealed 
in N2 ambient. From these results, they concluded the p-NiO did not act as the ISL to 
reduce the Schottky barrier height at the p-GaN/Au interface. 

Although, the resistivity of NiO can be decreased by an increase in Ni^"^ ions, or 
addition of monovalent atoms, such as lithium, or increased nickel vacancies and/or 
interstitial oxygen concentration [Ant92], NiO usually remain very resistive [Sat93]. 
Stoichiometric NiO is an insulator with a resistivity of the order of 10 Q-cm at room 
temperature [Adl70]. Because of the insulation of NiO, the poor contact performance in 
the work of [MaeOO] might be due to the large thickness of NiO film and non-epitaxial 
sputtered NiO in the contacts, resulting in larger resistance in the contact. The data from 
those contacts with thick NiO layers may be insufficient to evaluate the role of p-NiO, as 
proposed by Ho [Ho99b]. 

An interesting scheme was proposed to form ohmic contacts to p-GaN using 
internal electric field caused by polarization effects [LiOO]. Due to spontaneous and 
piezoelectric polarization effects, sheet charges can be induced in the AlxGai.xN/GaN 
materials system, and low specific contact resistance can be attained in contacts based on 
polarization fields at lower doping levels. In this way, a polarization-charge-based 
contact becomes a viable alternative to the formation of ohmic contacts to p-GaN. 
Experimental work showed that a spontaneous polarization vector could be formed and 
pointed to the substrate in the AlojGao.gN/GaN superlattice structure. Ni contacts to this 



CHAPTERS 
EXPERIMENTAL PROCEDURES 

3.1 Introduction 

This chapter describes the experimental procedures that were followed for sample 

cleaning and preparation of electrical contact to p-GaN. The procedures consisted of an 

initial cleaning of the samples followed by metal contact deposition. These contacts were 

then heat treated and characterized in terms of their electrical properties, surface 

composition, surface morphology and interfacial reaction products. 

3.2 Contact Preparation 
The p-GaN films used in these experiments were either purchased from SVT 
Associates grown with a Molecular Beam Epitaxy technique (referred to as MBE-GaN, 
1 (am-thick) or grown by low pressure Metalorganic Chemical Vapor Deposition 
(referred to as MOCVD-GaN, 2.5nm-thick). Mg and Si were used as doping elements in 
p- and n-type GaN, respectively. The SVT l^m-thick p-GaN epitaxy fihn was grown on 
the c-plane of sapphire substrates with an rf N2 plasma and a solid Ga source. The 
substrate temperature was of 700°C and the growth rate was of 0.5^m/hr. A 350A-thick 
AIN was used as a buffer layer before the deposition of a l^m thick p-type GaN film 
using elemental Mg as the p-type dopant. The GaN film was as-grown p-type with no 
post-growth activation. SIMS data showed about 10^° cm"^ for the actual level of Mg in 
the film. A hole concentration (p-type) of 1.1 -3.5 x lO'^ cm"^, and an electron 



46 



47 

concentration of 7 ~ 9 x lO'^ cm'^ (n-type) were measured by Hall measurement at room 
temperature for the MBE- and MOCVD- GaN. 

For the H2O2 treatment experiments, samples (5 mm x 5 mm) were cleaned with 
ultrasonically agitated acetone (5 min), methanol (5 min) and boiling aqua regia (10 min) 
sequentially before being flushed with DI water. All samples were blown dry with N2 gas 
between each step. These samples were then quickly placed into an electron beam 
evaporator system for contact deposition (using the "NOG" metallization principles as 
developed in Chapter 5), through a van der Pauw shadow mask [10] at a base pressure of 
<10'^ Torr. One sample was not treated further, i.e. kept in the as-cleaned state with 
contacts deposited. Other samples (with contacts on the surface) were immersed in 
H2O2/H2O solutions (1 : 1 and 1 :5) for 30 sec or 300 sec followed by DI water rinse, N2 
blow dry, and a short time (<60sec) dry at 80°C in air. 

The contact metals used in the "NOG" scheme consisted of Ni/Au, Ni/Ti/Au, 
Ag/Ti/Au, Ni/Al/Au, Pt/Au, Pt/Si/Pt/Au and Pt/Mg/Pt/Au with the first listed metal being 
deposited first and adjacent to the GaN. The Ni/Au and Pt/Au were mainly used for 
comparison, and the first layer of either gallide-forming metal (Ni, Pt), or neutral metal 
(Ag) was used to study the effects of different gallide-forming metals on the contact 
resistance. Ti, Al, Mg, Si were used as the second layer because they are nitride-forming 
metals. It is difficult for Al to absorb H from the GaN lattice, therefore metallization with 
Al were evaluated to test whether the contact resistance was decreased when using an Al 
layer. If it is, this would argue against H extraction postulated by Suzuki et al [Suz99]. Al 
reacts readily with O to form an insulator, AI2O3, which is expected to increase the 
specific contact resistance. Mg was used because it is the p-type doping element in p- 



48 

GaN. It was postulated that Mg in the metalHzation would increase the hole concentration 
in the contact interface region. Si was tested because of its strong tendency to form 
nitrides (Deh93). 

All samples were degreased prior to deposition using ultrasonicated acetone 
followed by methanol, each for 5min., and blown dry by N2. Any native oxide was then 
removed using boiling aqua regia for lOmin followed by a 5min DI rinse and N2 blow 
dry. Either regular (square) TLM patterns or circular TLM (CTLM) patterns, discussed in 
Chapter 2, was used for contact resistivity measurements. For square TLM patterns, the 
leakage path was isolated with a technique of inductively coupled plasma (ICP) dry 
etching using Ar + CI2 + N2 plasma under a pressure of 5 mTorr. For the CTLM patterns, 
the outer ring contact of the mask has a diameter (2 R) of 350/mi and the irmer dot 
contact's diameters vary (2 r) between 340 to 310 /^m to result in contact distances of 5, 
10, 15 and 20//m. Ohmic contacts are characterized by plotting the total resistance versus 
the contact spacings. The specific contact resistance and sheet resistance were derived 
from this plot as discussed in Chapter 2. The CTLM pattern was transferred to the 
multilayer metal contacts by a photoHthographic lift-off process. Positive photoresist 
(AZ1529) were used in these experiments. Prior to metal deposition the samples were 
etched with 10% diluted HF for 30sec to remove native oxides. The samples were then 
immediately introduced into the vacuum chamber for contact deposition. 

For contacts used in Hall measurements, the Van der Pauw configuration [Van58] 
was employed to determine both the Hall coefficient and the resistivity of the GaN films, 
as shown in Figure 3.1. Dot contacts of multiple metal layers of 25A Ni/500A Ti/500A 
Au and lOOA Ni/500A Ti/500A Au were used to form the ohmic contacts to p-GaN. 



49 




Figure 3.1 Configuration of contacts used in Hall measurement. 



These contacts were deposited through a shadow mask with the dot size of WOAmm, and 
the distance between two contacts of 4mm. 

All contacts were deposited in an electron beam evaporation system with a glass 
bell jar. The system was pumped with a Varian oil diffusion pump backed by a two-stage 
mechanical rotary vane pump, providing a base pressure of 10'^ ~ low 10"^ Torr. The 
metal charge in the electron beam well consisted of metal pellets with the following 
purities: Ni (Target Materials Inc., 99.98%); Ti (Cerac, 99.95%); Au (Materials Research 
Corporation, 99.95%); Pt(99.95%, Johnson Matthey, Inc.), Al (Cerac, 99.99%), Mg 
(Cerac, 99.99%), Si (semiconductor grade) and Ag (Cerac, 99.99%). The metal layer 
thickness was monitored using a quartz crystal oscillator. The contact thickness varied 
with different experiments and will be noted in the results sections. 

For RTA (Rapid Thermal Annealing) experiments, all metallization schemes were 
heat treated in a custom 50cm quartz tube with a 25cm hot zone and flowing N2 or O2 



50 

(high purity, 99.995%) as the ambient. The gas flow rate was monitored with a 
MANOSTAT® flow meter with a typical setting of 1 10 standard cubic centimeter per 
minute (seem) . • 

3.3 Characterization 

Contacts on GaN were characterized in the as-deposited state and following each 
of the above described heat treatments. Their electrical properties (I-V) and surface 
composition, surface morphology and interfacial reaction products were also 
characterized. A group of characterization techniques of I-V, light transmittance 
measurement. Hall measurement. Auger electron spectroscopy (AES) [Bru92], X-ray 
photoelectron spectroscopy (XPS) [Sib96], scanning electron microscopy (SEM) [Lee93] 
and transmission electron microscopy (TEM) were used in this study. 

The electrical properties of all contacts were investigated using room temperature 
current-voltage (I-V) measurements between two front surface dot contacts or the 
concentric dot/ring pattern described above. The I-V data were obtained by measuring the 
current flow between two adjacent top contacts under an applied bias. The ohmic or 
rectifying nature of the contacts could be determined by the linearity of the I-V curves 
and total resistance. The reverse-bias breakdown voltage of rectifying contacts could also 
be determined from their I-V characteristics. 

To evaluate the possibility of increased electrical conductivity from surface 
leakage current after the H2O2 treatment in Chapter 4, all of these samples were measured 
after contact deposition, and then cleaned with acid (IHCl + 2H2O) followed by base 
(IKOH + 2H2O) solutions for 10 min in each step, and rinsed with DI water after each 



51 

cleaning step. The I-V data were collected again without additional treatment. As shown 
below, the I-V data were unchanged by this treatment. 

The light-transmission characteristics of the Ni/Au and Ni/Au/Ni contacts (in 
Chapter 5) were measured with a commercial Zeiss UV grating monochrometer over the 
wavelength range of 300 to 700nm. The light transmission through the GaN film with 
and without the metal contact were measured on each sample at either a fixed or variable 
wavelength (Figure 3.2). The ratio of these two transmission values (I1/I2) at X, = 450imi 
was reported as the light transmittance through the Ni/Au "transparent" ohmic contacts. 
The repeatability of the data were checked with three sets of data obained fi^om each 
sample. The data reported are generally the average of these three measurements. 

For Hall measurements, a computer controlled MMR® commercial measurement 
system was used. The Van der Pauw method [Van58] was employed to determine both 
the Hall coefficient and the resistivity of the films. Dot contacts of lOOA Ni/500A Ti/ 
500A Au were used for p-GaN. All were deposited with an electron beam evaporator 



Metal contact 




Figure 3.2 Schematic of light transmittance measurement 



52 

through a shadow mask. The dot diameter was of 0.4 mm, and the distance between two 
contacts was 4 mm. A magnetic field of 3 kilo-Gauss was chosen automatically by the 
computer program to improve the accuracy of the Hall coefficient and resistivity values. 

Scanning electron microscopy (SEM, JEM 6400) was used in the secondary 
eletron or back scattering eletron mode to characterize the surface morphology of the 
grown films as well as the deposited layers in an attempt to determine the microstructure 
evolution and possible flaws that could cause high contact resistance. 

Transmission electron microscopy (TEM, JEOL 200CX) was performed with 200 
keV acceleration voltage for plane view analysis after H2O2 solution treatment to detect 
any defects in MBE-GaN 

The surface of the GaN film was characterized by atomic force microscopy using 
a Nanoscope III system fi-om the Digital Instruments, Inc. The AFM was operated in the 
tapping mode, and the height data was used in this work. 

Auger electron spectroscopy (AES) [Bru92] was used to measure the elemental 
composition of the atoms in the surface region using a Perkin-Elmer PHI Model 660 
scanning Auger microprobe (SAM). AES surface N(E) survey spectra from these GaN 
samples were recorded over the energy range of 50 to 2050 eV using a 5 keV, 30 nA 
electron beam with a diameter of ~1 |am. Depth profiles were also collected. With this 
technique, all elements of the layers can be determined except for H and He and generally 
down to a value of 0.1- lat%. The interfaces of the layers can be analyzed with a 
resolution approaching lOOA to determine possible compound formation. 

X-ray photoelectron spectroscopy (XPS) [Sib96] or electron spectroscopy for 
chemical analysis (ESCA) is an analytical technique similar to AES except the incident 



53 

energetical beam is x-rays. This technique is usually used for chemical state identification 
of surface species. XPS was used in this study to provide information on chemical 
bonding and compound formation that were not apparent from AES profiles. These 
measurements were performed with a Perkin-Elmer PHI Model 5100 ESCA system. Two 
large samples were used (10x10 mm) in this analysis. The ESCA data were collected 
using an Al X-ray source (Ka, hv = 1486.6 eV) and a hemispherical analyzer set at a pass 
energy of 35.75 eV. For these studies the ESCA source area was 4 mm x 6 mm. 

Secondary ion mass spectrometry (SMS) was used to determine impurity levels 
at the surface or in the films through depth profiling and was used to measure trace 
dopant profiles through the structure. SIMS is capable of quantifying the impurity levels 
for many elements depending on their detection sensitivity. However, SIMS has a limited 
possibility of providing useful information following heat treatments when diffusion 
distance are large. In this work, SIMS was used to detect the oxygen and hydrogen 
impurity levels in the H2O2 treated samples. This analysis were performed using a Cs"" 
primary ion beam and negative secondary ion detection. The raster size was 250 x 250 
jom^, and the counts were taken only from the central 75 x 75 |im . 



CHAPTER 4 
EFFECTS OF H2O2 SOLUTION TREATMENT ON p-GaN 



4.1 Introduction 

As described in Chapter 2, the high vapor pressure difference between gallium 
and nitrogen in GaN could lead to preferential loss of nitrogen and a gallium-rich surface. 
In this Chapter, stabilization of the surface of MBE grown GaN with H2O2 solutions is 
demonstrated. It is found that the H2O2 treatments are able to increase the carrier 
concentration by a factor of two. and lead to higher current levels through contacts for 
MBEp-GaN. 

As discussed in Chapter 2, oxygen is reported to react with GaN to form 
monoclinic beta-Ga203, with many polytypes being observed [Wol97, Roy52]. Annealing 
GaN in a O2 ambient is reported to dissociate the Mg-H complex and reactivate the Mg 
acceptors [Suz99]. Also, cleaning in common acids or bases or solvent is effective in 
reducing the amount of surface oxides and other contamination. The effects of aqua regia 
(IHNO3 + 3HC1) [Kim98], HCl [Kin96], HF [Kin96], KOH [Lee99], and (NH4)2S 
[CaoOO] on removing of native oxides on GaN have been reported. 

This chapter describes how H2O2 solution was used to clean/passivate the p-GaN 
film after contact deposition. Hydrogen peroxide was shown to increase the epilayer 
conductivity. Hydrogen peroxide is an active oxidant, and may be expected to react with 
GaN to form gallium oxide and/or gallium hydroxide to break up the Mg-H complexes 
and reactivate the Mg acceptors. These effects will be discussed. 

54 



55 

For I-V and Hall measurements, two types of metallization were used. For the 
treatments using 1H202:5H20 (volume ratios hereafter referred as 1:5), and 1:1 solutions, 
a contact scheme of 25 A Ni/500A Ti/500A Au was used, where the first metal layer is 
deposited onto the GaN epilayer. For treatments with 5:1 and "pure" (37%) H2O2, the 
contacts are 100 A Ni/500A Ti/500A Au, because this contact scheme led to higher 
current and more linear I-V curves, as discussed in Chapter 5. 

4.2 Modification of Electrical Conductivity 
4.2.1 Effects of H7 O 2 Concentration 

Remember, for these experiments, dot contacts were deposited onto GaN epilayer 
as described in Chapter 3 using a shadow mask., then the entire sample, including the 
contacts were treated with H2O2 solution. The I-V results after immersing the sample in 
1:5 or 1:1 solutions are shown in Figure 5.1. The current transported in the samples was 
found to increase as the soaking time and H2O2 concentration increased from 30 to 300 
sec and from 1:5 to 1:1, respectively. Compared to the as-cleaned state, an increase of 
~100% in the magnitude of current was found after immersion in the 1H202:1H20 
solution for 300sec as shown in Figure 4.1a. 

The I-V curves after immersion in the 5:1 and "pure" H2O2 treatment are shown 
in Figure 4.2. Straight, linear I-V curves were found for as-deposited contacts of lOOA 
Ni/500A Ti/500A Au treated with concentrated H2O2 solution (Figure 4.2), as compared 
to those with the 25A Ni/500A Ti/500A Au contacts and treated with a solution with 
lower H2O2 concentration (Figure 4.1). The highest currents were achieved in the 5:1 
treated samples. In the as-deposited state, after a "pure" H2O2 treatment, the highest 







56 


' 


Rf\ 










ou - 










40 - 


(a) 




' 




30 - 




^^ 


1 




20- 




^^^^^ 


' 




lio- 




^^^^ 






Current, 

o o 

1 1 


.^ 


^^"^ 






-20 - 
-30 \ 


^ 


P^ 








♦ As cleaned 




-40 - 






n1H2O2:5H2Ofor30sec 






-J3U \ 1 r 


1 1 1 1 r 1 1 




f 


-5 -4 -: 


3-2-1012345 








Voltaj 


38, Volt 





UVJ - 


(b) 




40 - 




^^ 


< 20- 




'^^^ 


k. 

5 -20 


^^ 






O As cleaned 


-40 t 


^^^ 


DlH2O2:1H2Ofor30sec 


i 


1 1 1 1 1 


A1H2O2:1H2Ofor300sec 


-60 - 


■-1 -r r 1 — ■ 



-5 -4 



-2-1012345 
Voltage, Volt 



? "..' 



Figure 4.1 Effects of H2O2 solution treatment on the I-V curves of 25 A Ni/500A 
Ti/500A Au to MBE p-GaN. (a) 1:5 solution; (b) 1:1 solution 



57 



zuu - 


(a) 




150- 






100- 


' •■■ 




< 50 

a 
♦J 

1 oj 

O -50 < 


^ 




-100^ 


As cleaned 


-150- 

[ 


1 1 r- ■' -I 1 


D 5H202+1 H20 for 30sec 
A 5H202+1 H20 for 300sec 


-200- 


— — I 1 1 1 



-5 A 



-2-1012345 
Voltage, Volt 



100 



50 



< 

=L 

♦J 

c 

I 

3 

o 







-50: 



-100 




-5 -4 



♦ As cleaned 

o pure H202 for 30sec 

A pure h2o2 for 300sec 



-2-1012345 
Voltage, Volt 



Figure 4.2 Effects of H2O2 solution treatment on the I-V curves of lOOA Ni/500A 
Ti/500A Au to MBE p-GaN. (a) 5:1 solution; (b) "pure" (37%) H2O2 



58 




50 100 150 200 250 300 350 
Immersition Time (sec) 



Figure 4.3 Hall measurement results of MBE p-GaN after immersion in 1:1 
solution. 



current increase after 300sec was -50%, while the increase after a 5:1 solution treatment 
was -200%. For samples treated in 1:5 and 1:1 solutions, the increases in current at 5V 
are 41% and 89% respectively with a treatment of BOOsec. 

The Hall data for the samples treated with 1:1 solutions are shown in Figure 4.3, 
An increase of about 70% was found for the carrier concentration after immersion for 
300sec, consistent with the increased current found for the treatment. The decrease in 
hole mobility is within experimental errors. 



4.2.2 Effects of Extended Immersion Time 

Using the 5:1 solution, the effects of 20 to 60 min extended treatment times were 
Studied. The I-V data are shown in Figure 4.4 and the Hall data are shown in Figure 



59 



200 



< 

♦J" 

c 
o 

k. 
k. 

3 

o 





^i«««"«""* 



,„«UM««' 



o As cleaned 

D 5H202:1H20 45sec 

A5H202:1H20 5min 

X5H2O2:1H2Ofor20min 

• 5H2O2:1H2Ofor60min 



-5^-3-2-1012345 
Voltage, Volt 



Figure 4.4 Effects of H2O2 solution treatment with extended time on the I-V 
curves of lOOA Ni/500A Ti/500A Au to MBE p-GaN. 



4.5. Compared to as cleaned (no H2O2 treatment), an immersion time of 20min resulted in 
a 200% increase in the current level. As the immersion time was further increased to 
60min, the current level dropped to -50% of the as cleaned state.These Hall data show 
that the hole concentration and mobility first increase and then decrease as the immersion 
time increase, consistent with the I-V data. 

For MBE-GaN, n-type GaN was also treated with H2O2 for times of 10, 30, 300, 
1200 and 3600sec using the 5:1 solution. The Hall data are shown in Figure 4.6. Short 
treatment times did not change the electron concentration or mobility within experimental 
error. At 30 and 300sec, the electron concentration increased (-12%), but these changes 
and the mobility changes are still in experimental errors (-15%). 



60 




10 



100 1000 

Immersition Time (sec) 



10000 



Figure 4.5 Effects of 5H2O2: IH2O solution treatment with extended time on Hall 
measurement results of 1 OOA Ni/500A Ti/500A Au to MBE p-GaN. 



All these results were obtained on epilayers of p- or n-GaN grown on sapphire by 
MBE, as described in Chapter 3. For comparison, epilayers grown by MOCVD were also 
treated with the same solutions and methods. The I-V and Hall data as shown in Figure 
4.7, and no changes above the experimental noise levels were detected, even after 
treatment time up to 60min. Thus the effects of peroxide treatment depend upon the 
growth history of the p-GaN 



4.2.3 Stability of the Increased Electrical Conductivity 

The stability with time of these modified electrical properties was monitored with 
I-V and Hall measurements every two days for two weeks. The improved electrical 



61 




10 100 1000 

Immersition Time (sec) 



10000 



Figure 4.6 Hall measurement results of 5H2O2: IH2O solution treated MBE n- 
GaN 




10 100 1000 

Immersion Time (sec) 



10000 



Figure 4.7 Effects of H2O2 treatment on MOCVD p-GaN. 



62 



E 
u 



10 



8- 



c b 
o 
'•^ 
CO 

it 

c 

o ^ 
c 
o 
O 

o 

i 2 



04 



(a) 



D- 



-n — Hole Concentration 
-e — Mobility 




t 6 8 10 

Immersion Time, day 



12 



15 



13 



o 

tt 

M 

11 5: 

E 
u 

9 t 

o 



-7 



14 



200 




4 6 8 

Immersion Time, day 



Figure 4.8 Hall measurement on stability of H2O2 treated samples, (a) n-GaN; (b) 
p-GaN. Both were treated with 5:1 solution for 20min. 






63 



Table 4.1 Atomic concentration of elements from AES surface survey analysis in MBE p- 
GaN after a xH202:yH20 treatment for either 30 or 300sec 



As cleaned 1 :5 for 30sec 1 :5 for 300sec 1 :1 for 30sec 1:1 for300sec 



C 4.9 6.3 5.2 5.9 3.5 

d 1.9 1.0 1.1 " 0.8 1.0 

Ga 52.0 48.5 47.6 47.2 49.0 

N 33.7 36.4 39.1 37.5 39.4 

O 7.6 7.8 7.0 6.7 7_A 



conductivity was found to be very stable. Figure 4.8 shows the Hall data for (a) p-type, 
and (b) n-type GaN treated with 5:1 solution for 5min and 20min respectively. 

4.3 Structural Characterization 
4.3.1 AES . 

Auger peaks from Ga, N and O plus C and CI were detected from samples after 
aqua regia cleaning and minimum exposure to air. The surface composition changed after 
the H2O2 solution treatment, as shown in table 4.1. Compared to the as-cleaned state, a 
slight increase of N (3 ~ 6%) and a slight decrease of Ga (3 ~ 5%) were found. Small 
increases of C (except for the 1:1 for 300sec treatment) and decreased CI also resulted 
from the H2O2 soaking. The atomic concentrations of oxygen remained fairly constant 
instead of increasing with the H2O2 immersion. This result was surprising because oxide 
or hydroxide formation was expected from this treatment of the GaN surface due to the 
reaction with H2O2. 



64 



4.3.2 ESCA 

The ESCA spectra for the Ga 2p, Ga 3d and 01 s photoelectron peaks are shown 

in Figure 4.9 for the as-cleaned (boiling aqua regia, no H2O2 treatment) and soaked 
(1H202:1 H2O solution for 300sec) surfaces. The Ga2p"^ and Ga2p^'^ peaks are at 
binding energies of 1 1 18.5 and 1 145.4 eV, similar to standard data [Mou95]. Comparing 
the peaks of samples before and after H2O2 immersion, no binding energy shift in the 
Ga2p"^ and Ga 2p^''^ spectra (Figure 4.9-a) were found. However, energy shifts of 0.85 
and 1.1 eV to higher binding energies were measured for the Ga 3d (Figure 4.9-b) and O 
Is (figure Figure 4.9-c), respectively, as shown in Table 4.2. 

4.3.3 SIMS 

Secondary ion mass spectrometry (SIMS) depth profiling was performed to 
determine any changes in the Ga, N, O and H levels. Two different sites were analyzed in 
each sample. While no significant variation in Ga and N levels were found, large 
differences in the levels of H and O were found between the two sites on each samples. 
The results are shown in Figure 4.10 for each sample. These results suggested that the H 
and O levels were not uniform and had large variations even over the same sample. 
Because the impurities were not uniform, systematic comparison of O and H levels with 
the change of immersion time was impossible by SIMS. 

4.3 A AFM 

The AFM technique was used to characterize the surface morphology of both 
MBE- and MOCVD-GaN epilayers with and without H2O2 treatments. 



65 



15 



3 



<A 

c 

4) 



(a) 




1:1 for 
300sec 



1150 



1140 1130 1120 

Binding Energy, eV 



26 



24 



22 20 18 

Binding Energy, eV 



16 



545 540 535 530 

Binding Energy, eV 



525 



>. * 



1110 





520 



Figure 4.9 Comparison of XPS peaks from as-cleaned and 1:1, 300sec H?02 
treated GaN. (a)Ga 2p; (b) Ga 3d; (c)0 1 s. 



66 



Table 4.2 XPS results from a 1:1, 300sec H2O2 cleaned p-GaN sample 



Surface Treatment 



Ga2p3/2 Ga2pl/2 Ga3d 01s 



As-cleaned 1118.5 1145.3 19.4 532.9 

lH2O2:lH2Ofor300sec 1118.5 1145.3 20.25 534.0 




200 400 600 800 

Sputter Time (sec) 



1000 




200 400 600 

Sputter Time, sec 



800 



1000 



4 A ^> 



Figure 4.10 Negative SIMS depth profile for 5:1 H2O2 solution treated p-GaN 

film (a) as-cleaned state; (b) immersed for 45sec; (c) immersed for 5min; 
(d) immersed for 20min and (e) immersed for 60min. Data from 2 
different points (a, b) on each sample are shown 



67 








200 400 600 800 

Sputter Time, sec 



1000 








200 400 600 800 

Sputter Time, sec 



1000 












^^^H1 \ bHI 



200 400 600 800 

Sputter Time, sec 



1000 



Figure 4.10-Continued 



68 



500 



NP 




1.00 



0.75 



0.50 




0.25 



Figure 4.1 1 AFM images of GaN. (a) MBE-GaN, RMS roughness is of 2.6nm; (b) 
MOCVD-GaN, RMS roughness is O.Vnm.The points on the MBE-GaN 
surface labeled NP are potentially nanopipes as discussed in the text. 



69 



Defects 




Figure 4.12 Microsturcture of MBE-GaN (TEM, plane-view) 



The AFM images for both types of samples are shown in Figure 4. 1 1 . The as- 
cleaned MBE-GaN and MOCVD-GaN had RMS surface roughness of 2.6 run and 0.7 
nm, respectively. Not only is the MBE epilayer rougher, but they also exhibited 
numerous areas labeled "NP" on Figure 4.1 la. As discussed below, these are possibly 
nanopipes which allow quick transport of atoms throughout the epilayers. The density of 
the "NP" points was ~10'° cm"^. Microstructure obtained with TEM plane view 
observation of the MBE GaN showed the presence of defects, as in Figure 4.12. While 
the character of the defects were not studied in detail, the larger dark defects are 



70 




Voltage, Volt 



c 
o 

3 

o 



0.01 -1 


(b) 


Jf 


0.006 - 




r/ 


0.002 




^y^ 


0.002 - 


/f^ 






■0.006 - 


—0— Untreated 






// 


-0-5:1 foreOsec 




-0.01 - 


-I 1 





-1 -0.6 -0.2 0.2 0.6 1 

Voltage, Volt 

Figure 4.13 Effects of 5:1, 60sec H2O2 treatment on the I-V of 500A Ni/500A Au 
to MBE p-GaN. (b) is the same to (a) but in a smaller scale. 



consistent with nanopipes. Their density, while still large, is only 10^ cm" compared to 
the density of 1 0'° cm"^ reported from Figure 4. 1 1 -(a). 



-. -^^ 



71 

4.4. Application in Formation of Ohmic Conta ct to p-GaN 
As shown above, the peroxide treatment of MBE p-type GaN can increase the 
hole concentration although the immersion time and solution concentration must be 
controlled.Increased hole concentration would allow formation of a lower resistance 
ohmic contact, so this technique was combined with the Ni/Au contact scheme for both 
MBE and MOCVD p-GaN. Figure 4.13-(a) shows the I-V data for Ni/Au contacts on 
MBE p-GaN. hicreased currents at the same voltage was found for voltages > 2V after 
treatment with a 5:1 solution for 60sec. However, at voltages < 1 V, the current was 
lower than for an untreated surface, as shown in Figure 4.13-(b). Lower currents are 
attributed to incomplete removal of oxide from the contact interface. The results from the 
MOCVD p-GaN show similar results. As a consquence, the increased conductivity of 
MBE-GaN did not result in reduced contact resistance as expected, presumably due to the 
dominance of interfacial reaction layers on current transport. 

4.5 Discussion 
Considering both hole and electron carriers, the electrical conductivity of a 
semiconductor material is usually described by [May90] : 

C7 = q-{n-fi„+p-Mp) (4-1) 

where a is electrical conductivity, q is the electron charge, n and/> and //„ and //p are 
electron and hole concentrations and mobilities, respectively, hi a p-type semiconductor, 
the majority carriers are holes, so the effects of electrons could be omitted because their 
concentration is so small. The electrical conductivity is only related to the hole 
concentration and mobility. The opposite is true for the case of n-GaN. From Figure 4.3, 



72 

the hole mobiUty is shown to be constant within experimental error, but the carrier 
concentration increased after immersion in the 1H202:1H20 solution. 

As the Hall contacts were deposited before the treatment in H2O2 solutions, the 
effects of H2O2 on the metal/GaN interface can be omitted. Possible reasons to account 
for the increased electrical conductivity after the H202:H20 immersions then might 
include: (a) a surface layer (e.g. gallium oxide) could form which if more conductive than 
GaN could lead to surface leakage currents, (b) etching of GaN by the H2O2 solution to 
reduce the epilayer thickness, (c) reduced hydrogen passivation of Mg, or (d) formation 
of Ga-0 compounds to stabilize the excess Ga atoms and decrease the concentration of N 
vacancy donors. 

A reaction of the type 

yH^O^+xGa =Ga^Oy+ yH^O . (4.2) 

was possible in this experiment, resulting in a GaxOy oxide. As no data is available to 
calculate the gallium oxide thickness in this reaction, the thickness was assumed to be 
very thin. If this oxide layer (or other possible reaction products) was more conductive 
than the GaN, then the increased conductivity after H2O2 immersion could be explained 
by a surface leakage current between the contact structures. However, literature data 
show that Ga203 is normally an insulating dielectric layer on the GaN [Ren99]. Various 
acid and base solutions [Sam73] were reported to be effective for removal of the Ga203 
layer. 

If the peroxide treatment created a conductive surface layer, then the conductivity 

should decrease after re-cleaning the surface with an acid or base solution. In contrast, the 

. ■ -'■> 

■■,- •' . .V 



73 



4000 



1-10 




110'' . 1.10 

Carrier concentration, cm^-3 



110 



Vbi=leV 

---- Vbi = 2eV 

- - Vbi = 3eV 



Figure 4.14 Relation among width of depletion region, carrier concentration and 
built-in potential in GaN 



I-V data after the acid and base solution cleaning showed a slightly increased (rather than 
a decreased) conductivity, so surface leakage current is discounted. 

If this Ga203 layer can create a depletion region in p-GaN due to the band 
bending effects, the carrier concentration might be changed. To consider this effect, the 
depletion width in p-GaN was calculated. The equation used in this calculation is 
[Sze81]: ' v 



^ w- ''^^^^'■ 



qN^ 



(4.3) 



74 

where W is the depletion width, £ = £s Spis semiconductor permitivity, ^ is 
semiconductor dielectric constant, So is vacuum permitivity, Fi,, is buih in potential, A'^ is 
acceptor concentration, and q is the magnitude of electronic charge. Using the value of e^ 
= 10.4 for GaN [Ho99b] and ^ = 1.6 x 10"'^ C, the width of the depletion region in GaN 
related with carrier concentration was calculated for different built-in potentials as shown 
in Figure 4.14. At the hole concentration of 2-3 x lO'^ cm'\ the depletion region is 
-1000 A. Compared to the thickness of 1 ^m, this depletion region is narrow, and is not 
expected to cause carrier change of 50% or higher. 

Oxygen impurities incorporated into p-GaN from the H202:H20 solution could 
cause a surface conversion to n-GaN and result in increased I-V conductivity. When 
was implanted into GaN and aimealed at 1 100°C [Abe96], it created n-type doping with 
an ionization level of -29 meV. Seifert et al [Sei83] proposed that oxygen 
substitutionally incorporated into nitrogen sites could be the origin of the free electron 
carriers in highly conductive n-type GaN. fri the present case, oxygen incorporation 
should increase the electron concentrations in n-GaN, but decrease the hole concenfration 
in p-GaN because of compensation. This might be used to explain any increased 
electrical conductivity in n-GaN, but not in p-GaN. Thus, an increased oxygen level from 
H2O2 treatment is also discounted. 

The etching of GaN can also lead to the change of carrier concentration in the 
measured results. Pearton et al [Pea93] have reported on the chemical etching of GaN 
films in aqueous 30-50% NaOH solutions at an etch rate of -2mn/min, although other 
authors [You97, Min96, Var96b] did not observe any open-circuit etching in the dark in 
alkaline solutions. Weyher et al [Wey97] found the etch rate of [0001] oriented bulk GaN 



75 



crystals and ep.taxial GaN layers .n aqueous KOH to depend upon the face exposed to 
the etching solutton. The N-tenninated face ,s etched in alkahne solutions bu, the Ga- 
te^nated face remains unaffected. Stmilar crystallog^phte dependency of wet etchtng 
«as also reported by Stocker et al [Sto98]. I. add.tion to open-crcuit etching in me dark, 
studtes were also perfonned on open circuit photoetehmg of GaN. M.nsky et al lM,n961 
noted photoetchtng of GaN tn aqueous KOH and HCl solutions. The etch rate seemed to 
he an order of nta^itude higher .n alkal.ne soluttons. Extended study of the photoetching 
behavior of GaN tn KOH solutions [You97] found that the etch rate was dependent on 
the mcdent Ught tntenstty, the dop.ng of the material, and the KOH concentration (pH 
values of to solutions). Peng ., a, [Pen98a, Pen98b] also demons^ated the open-circuit 
photoetchtng m aqueous KOH and H3PO, solut.ons to be pH dependent.When 
illumtnated from a 253.7.nm mercury line source, etching of GaN was observed in 

. c 1 1 f„ 1 ^ The GaN etch rate was 90nit)/miii at pH 
solutions with pH values rangmg from 1 1 to 1 5. The GaN 

= 14.25, but drops rapidly for pH values above or below 14.25. 

The relation of pH value wtth the HA concentratton in the H.O,/H.O solutions 
,s shown in Kgure 4.15. For the solution used in this work, the pH value of the 5:1 
solution is around 5.5, of the 1:5 solution ts about 4.7 and of the -pure" H.O. solution is 
around 4.6 ~ 4.7 (Note the concentration of H,0, in "pure" H,0. ts about 37%.), which 
are significantly smaller than the reported peak values of 12 or over. 

If the GaN has been etched by the H,0. solution, the GaN film thickness would 
aecrease and the apparent carrier concentratton would decrease the ongtnal values of film 
thtckness was used. If only the etchtag effect was involved, the measured carrier 
concentration would decrease monotonically. Thts prediction ts incons.stent w.th the 



"T^'- 



76 




20 40 60 80 100 

H2O2 Concentration (%, w/w) 



Figure 4.15 Relation between pH values and H2O2 concentration 
in H2O2/H2O solutions [PerOO] 



results obtained in this study where the carrier concentration increased with short time 
treatment and decreased with longer time of treatment. Besides, the measurement of GaN 
epilayer thickness with profilometer and depth profiles showed no significant difference 
in the film thickness. 

Another possible explanation of increased electrical conductivity may be related 

to the reaction of 

OH-+H*=H,0 (4.4) 

i.e. the removal of H fi-om the GaN film. It is well known that H compensates Mg 
acceptors in p-GaN [Sug98, Ama89]. hi fact, a similar mechanism of H removal fi-om 
GaN was used to explain the low resistance of Ta/Ti contacts to p-GaN [Suz99]. It was 
speculated that Ta/Ti can withdraw (absorb) H from Mg-H complexes, thus reactivating 
the Mg acceptor near the contact interface. If this argument was correct, the I-V data 
would be expected to continually increase with extended treatment time because more 



77 



hydrogen could be removed and more Mg acceptors reactivated. The I-V data in this 
work showed the conductivity decreased as the immersion time was extended. So the H 
removal cannot explain all of the results obtained. Additional physical or chemical 
processes must be involved during the immersion of GaN epilayers in H2O2 solutions, as 
discussed below for surface injection of isoelectronic oxygen donors. 

Finally, formation of gallium oxides can reduce the excess Ga atoms, resulting in 
a decreased nitrogen vacancy concentration, and the hole concentration should increase 
due to less compensation by nitrogen vacancies, as discussed in Chapter 2. 
In addition to the formation of Ga oxides, dissociation of H2O2 to H2O to yield O could 
result in the formation of chemisorption bonds with GaN. This would explain the increase 
in binding energies for the Ga3d due to an increased net charge on the oxygen versus 
nitrogen atoms resulting from the larger electronegativity of oxygen [Deh93]. The 
oxygen electronegativity (3.44 on the Pauling scale) is larger than that of nitrogen (3.04), 
so higher gallium bonding energies were expected for the Ga-0 versus Ga-N bonds. 
Oxygen bonds decrease the electron concentration in the Ga valence orbitals which 
results in a weaker screening effect of the nuclear charges by the inner electrons, so the 
binding energy of Ga3d should increase as observed [Cza75]. If only GaxOy formed, 
these arguments predict a lower binding energy for the 01s, whereas an increased 
binding energy was measured. This suggests that Ga(0H)3 formed rather than GaxOy. The 
argument of Ga(0H)3 formation is consistent with increased binding energies of both the 
Ga3d and 01s in the results shown above. It also could be consistent with removal of H 
from the Mg acceptors to form hydroxyl bonds. The typical depths of hydrogen 
incorporation and maximum concentration measured in various processing steps were 



78 



Table 4.3 Relation of hydrogen incorporation and processing steps [Pea97b] 



Process 


Temp, 


"C 


Maximum [H] (cm"'*) 


Incorporation Depth (|im) 


H2O boil 


100 




10^^ 


1.0 


PECVD SiNx 


125 




.3x10'' 


0.6 


Dry Etch 


170 




10'^~10^° 


>0.2 


Implant isolation 


25 




Dose dependent 


2.0 


Wet etch 


85 




2x10'^ 


0.6 



compiled in Table 4.3 [Pea97b]. Based on these data, it is clear that even at room 
temperature, such large incorporation distances have not been demonstrated for nitrogen 
vacancies or changes in excess Ga profiles. 

A model of the amphoteric roles of O is presented here. One effect of O could be 
the combination of excess Ga and O (or OH') to form stable compounds with reduction of 
donor N vacancies, but a second effect could be the incorporation of donor O impurities 
into the GaN lattice which should compensate Mg acceptors. 

Formation of Ga(0H)3 could also modify the concentration of excess Ga and 
effectively reduce the N vacancies in the GaN film. The high vapor pressure of nitrogen 
above GaN could lead to desorption of N, excess Ga and high concentration of 
compensating donors in the surface region. Immersion of GaN samples in H2O2 solutions 
could potentially reduce the concentration of excess Ga atoms by formation of oxides or 
hydroxides. In the literature, oxidation of Ga with O leads to several polytypes for 
oxidation products [Wol97]. 

Theoretical [Neu995] and experimental [Wet97] studies have also shown that 
oxygen is a relatively shallow donor which can be present at high concentrations (lO"^ ~ 
10 cm" ) in undoped GaN films. High concentration of oxygen resulted in high 
background free electron concentrations commonly observed in as-grown films. The 



79 . , 

formation energy for substitution of O onto the N sites is much lower than onto the Ga 
sites, and the covalent radius and ionicity of O and N are similar. The group VI element 
O would contribute one extra electron to the conduction band and increase the free 
electron concentration after incorporation into the GaN lattice. The electrical conductivity 
should decrease for the p-GaN, as observed. In this experiment, the decrease of electrical 
conductivity after a long time (60min) immersion is believed to result from this 
mechanism. 

As the immersion time increased, two processes could compete to influence the 
change of carrier concentration. Reduction of compensation by removal of H, excess Ga 
or N vacancies should increase the hole concentration and conductivity. As more oxygen 
atoms were incorporated substitutionally into the GaN films decreased hole concentration 
and mobility would be expected at long times. 

While the competing effects at short versus long time on excess Ga and Vn 
qualitatively explain the experimental results, the question remains as to whether there is 
sufficient atomic mobility at room temperature to make this a viable mechanism. 

The difftision of oxygen in GaN has been studied using the '^O in Si02/GaN 
interface [Pea99] and the diffusion depth was found to be -400 ~ 500A for as-deposited 
films. It was also found dislocation pipe diffusion led to larger regions of diffusion 
around threading dislocations. The overall transport of the oxygen is dominated by the 
pipe diffusion which is along the dislocation axis only. The square of the diffusion 
distance in the dislocation is given by 

X^=2-Dj-t " (4.5) 






80 

where Da is the diffusivity in the dislocation, and t is the diffusion time. The least square 
fit to the data led to the equation [Pea99]: 

D = (4.5±2.2).10-'^-exp( ) (4.6) 

k-T 

where k is the Boltzmann's constant. The activation is approximately half of the expected 

value of diffusion in bulk material. 

Using these two equations (4.5 and 4.6), the diffusivity was calculated to be 
between 2.38 x 10''^ and 8.8x10"''' cm^/sec (using the upper and lower limit of the data in 
equation 4.6) and the diffusion distance at which the oxygen concentration equals to the 
background concentration of lO' ''cm'^ was between 18.5 A and 0.355)am for a time 
period of 60 min. With the experimental times used in this study, it is difficult to believe 
that the oxygen could diffuse deep enough into the GaN film at room temperature to react 
with excess Ga over a \\xm epilayer by a uniform transport mechanism. It is much more 
reasonable to expect rapid diffusion down to a defect perpendicular to the surface, such 
as a dislocation core or nanopipe, with diffusion out of the core or pipe, parallel to the 
surface, over much shorter distances. 

Defects like nanopipes, screw dislocations and others have been reported in GaN 
and are known to be paths of enhanced mobility for oxygen and hydrogen. Nanopipes 
with diameters ranging from 5nm to 0.5|im were reported at a densities of ~10 cm" . 
They were reported to be parallel to the c-axis of GaN unit cell [Ven99, Kan99] as shown 
in Figure 2.1. The main composition of nanopipes was Ga, C and O [Kan99], as 
discussed in Chapter 2. Considering the high chemical activity of O in the H2O2 solution, 
diffusion of O or H along these nanopipes and out of the GaN or nanopipes into the 



81 

surface or GaN lattice would be logical and would explain the results from H2O2 
treatments. With the help of these nanopipes, the point defects are expected to be 
extracted or to reach the deep site of the epilayer easily. . 

Based on this model of enhanced atomic diffusion along nanopipes, the expected 
increase of carrier concentration from passivation of these defects (using the nanopipe as 
an example) was predicted as following. 

For the perfect crystals with no nanopipe, the total concentration of holes would 



be 



N = N,-S-t 



(4.7) 



where A^, is the average hole concentration in this perfect crystal, S is the sample surface 
area, and t is the sample thickness. With the incorporation of nanopipes, the amount of 



c 
o 
u 

1 

U 




50 100 

Distance, y = (i'2 + X, nm 

■ Ni=l(r9cnf^-2 

■ Ni=l(ri0cnir-2 



150 



200 



Figure 4.16 Relation of carrier concentration decrease and nanopipe density in 
GaN 



82 

carriers will be changed. If the dislocation/nanopipes are perpendicular to the sample 
surface, extending through the thickness of epilayer film, and the hole carriers are 
completely removed in the dislocation/nanopipes and over the depletion distance around 
the dislocation/nanopipes without effecting the carrer density in the "perfect" region 
beyond the depletion distance, the reduced carriers concentration Nh' would be: 

N,'-S-t = N,-[\-7r-(i + xf-N^yS-t (4.8) 

where x is the average depletion distance around the nanopipes, d is the average nanopipe 
diameter, and A/p is the nanopipe density with a unit of cm" . 

The reduced carrier concentration compared to the perfect crystal will be: 

AN„=^ = l-7r-ii+xf-N^=\-7r-y'-N^ (4.9) 

where since x is of the same order of magnitude as d and exerts the same influence on the 
change of hole concentration, the values are added together as defined as^'. Figure 4.16 
shows the relationship between carrier concentration decrease and the influence of 
dislocation/nanopipe size and density. At a density A'p of 1 cm" , the carrier 
concentration would be reduced by 50% with ay distance of 126nm. At a nanopipe 
density of 10'° cm"^, the >' distance of only 40nm would result in the same reduction. 

In the AFM image shown in Figure 4. 1 1-a, the "cavities" labeled "NP" are 
suspected to be nanopipe or dislocation sites and their density is calculated to be 4x10 
cm'^ with a size of 10 ~50nm. The >' distance for p-GaN with our measured carrier 
concentration increase of -100% is calculated to be 40nm. These values are in good 
agreement with the calculated depletion of holes assuming the nanopipe surface cause 
depletion of carriers. The variation in O and H fi-om SIMS data (Figure 4.10) would be 



10 



consistent with the literature data [Kan99] showing that nanopipes were rich in Ga, O and 
C. The result that O and H at one point was continually higher than another one was 
consistent with the published results, and is also consistent with the assumption that the 
direction of these nanopipes were parallel to the c-axis of GaN unit cell. It should be 
noted that analysis size of SIMS is 75 x 75 f^m^ which means each analysis covered an 
area containing many nanopipes. The results from the SIMS analysis in this work suggest 
the nanopipe/dislocation sizes and/or density are not uniform. 

4.6 Summary 
The effects of H2O2 treatment on GaN had been studied using the I-V, Hall, AES, 
ESCA, SIMS, AFM and TEM. It was found that the electrical conductivity increases after 
short (< 20 min.) H2O2 solution treatments with contact already deposited. The 
magnitude of the current change was related to the H2O2 concentration, with 5H2O2: 
IH2O solutions giving best results. The Hall data showed that while the carrier mobility 
exhibited changes <10%, the average hole concentrations could be changed by up to 
100%. According to AES data, the Ga atomic concentration was decreased, N increased 
and O not changed after the H2O2 treatment. Using the ESCA analysis, formation of 
Ga(0H)3 was shown to result fi-om H2O2 treatments. With SIMS depth profiling, big 
variations of H and O were found across the samples, consistent with published work and 
AFM data from this study. Based on the possible reactions between GaN and H2O2, 
reasons for the increased conductivity of MBE grown p-GaN were discussed, and a 
model of reaction with oxygen to deplete nitrogen vacancies or compensating H at short 
time, and injection of O donors at long times, was proposed. 



CHAPTER 5 
"NOG" SCHEME FOR OHMIC CONTACT TO p-GaN 



5.1 Introduction 

When Maruska and Tietjen [Mar69] used a chemical vapor deposition technique 
to make GaN layer in the late 1960's, the GaN was highly conductive n-type even when 
not deliberately doped. Two dominant mechanisms were speculated to explain intrinsic n- 
type as-grown materials. One was nitrogen vacancies and another was of unintentional 
oxygen doping. The N vacancy was considered to be a donor because a N vacancy would 
form a void surrounded by four Ga atoms contributing three electrons. Two of these three 
electrons could reconstruct and leave a single electron that could be donated to the 
conduction band. This model was later questioned, and instead unintentional oxygen was 
proposed to be the donor in as-grown GaN [Pan73]. Oxygen with its six valence electrons 
on a N site (N has five valence electrons) would be a single donor. 

The proposed "Nitride-forming metal Over Gallide-forming metal" ("NOG") 
scheme in this chapter is based on interfacial reactions to control the Ga and N to 
decrease N vacancies. Based on the review above and in Chapter 2, N vacancies would 
be at least one of the reasons for a high concentration of donors in GaN 

Metallurgical reactions of transition metals with GaN have been reviewed in 
Chapter 2. Only the principles of the "NOG" scheme and its applications are presented in 
this chapter. Applications of "NOG" include interpreting published literature results and 
designing/testing new contact schemes in this study. 

84 



85 



5.2 Principles of "NOG" Scheme 

Since there is a large difference in the vapor pressures of Ga and N, there may be 
a high concentration of N vacancies, VN.in as-grown epilayers. A high Vn condition is 
equivalent to a Ga-rich condition. An opposite situation could be postulated: if a N-rich 
condition could be created in as-grown GaN films, which is equivalent to creating Ga 
vacancies, the as-grown GaN should be intrinsic p-type. This condition has not been 
achieved in bulk GaN films, probably because of the high vapor pressure of N in 
equilibrium with GaN. histead, this postulated condition leading to p-GaN might be 
achieved by interfacial reactions in the contact region. If extra N atoms could be kept 
between the contact metal layer and the bulk p-GaN film, a N-rich condition could be 
formed at the metal/GaN interface. The extra N atoms could fill the Vn positions and 
create Ga vacancies acceptors. If such Ga vacancy acceptors were shallow and reached a 
sufficient concentration (>10'^ cm"^), the interfacial region could become p"^-GaN and 
current transport could be dominated by field emission. A low resistance ohmic contact 
could be obtained as a result. Even a decrease in the Vn concentration should increase the 
fi-ee hole concentration in the contact region. These are the postulates upon which the 
"NOG" scheme is based. While these reasons seem sound, the "NOG" scheme has not 
led to new and improved contacts. It is worth a while however to review the progress 
made using these ideas. 



86 



Protection metal (Au) 



Nitride-forming Metal 



Gallide-forming metal 



p-GaN 



Figure 5.1 Principle of "NOG" scheme. 

According to the enthalpy of the metallurgical reactions, all transition metals were 
classified into three groups in Chapter 2: gallide-forming, nitride-forming and neutral 
metals with respect to reactions with GaN. Metals for forming "NOG" contacts were 
selected based upon this classification. 

The structure of a "NOG" contact is illustrated in Figure 5.1. A gallide-forming 
metal adjacent to GaN is followed by a nitride- forming metal, which is covered with a 
layer of protective metal (such as Au). Under suitable annealing conditions, the gallide- 
forming metal reacts with GaN to form stable gallides and release N atoms. This first 
metal layer must both dissociate the GaN lattice and prevent or slow down the nitrogen 
out-difflision. The nitride-forming metal layer would help to keep the released N atoms in 
the contact interfacial region and create a high N chemical potential. 



5.3 Comparison with Published Contact Results 
The achievement of low resistance ohmic contact to p-GaN is of great importance 
to GaN device performance. Many studies have been reported along with postulated 



87 

mechanisms to explain reduced contact resistance, such as interfacial reactions 
ehminating barriers [Hol97] and doping the surface region [Tre96, Tre97], GaN re- 
growth [You98, Tre96], H extraction [Suz99] and Ni oxidation [Ho99a, Ho99b]. A few 
representative contact schemes are discussed here based on the principle of "NOG" to 
show the possible appUcations of this contact scheme. The discussion remains strictly 
speculative. 

The contact scheme of Ni/Au [Fun99, Kim97a, Ish97, Kin97, Tre97] has been 
widely used for GaN device fabrication. Based on the "NOG" principles, this lower 
resistance for these contacts results from the reaction between Ni and GaN[Ber93, 
Guo96, Ven97]. The Ga would react with Ni to form stable gallides and reduce excess Ga 
atoms. Reduced excess Ga is expected to result in reduced concentrations of Vn and less 
compensation of acceptors. This would result in higher free hole concentrations in the 
interfacial region. The success of some other contact schemes, like Pd/Au[Kim98, 
Kim97b], Pt/Au [Kin97, Mor96, Yoo96], Pd/Au/Pt/Au [Kim97b], Pd/Pt/Au [Kin97] and 
Pt/Ni/Au [Jan99], could be explained by the same mechanism. As discussed in Chapter 2, 
Pd and Pt formed more stable gallides than Ni, so it is not surprising that contacts with 
lower specific resistance were obtained with Pd/Au and Pt/Au contacts 

A relatively low specific contact resistance (3.6x1 0''' Q-cm^) was obtained with 
Ni/Zn-Au [You98] on p-GaN with a carrier concentration of Nh = 4.4x1 o'^cm"^. The 
authors reduced the time between p-GaN film growth and contact metal evaporation in a 
high vacuum system. They postulated that Zn was an acceptor and that the Zn-Au alloy 
layer increased the interface carrier concentration. Zn is an acceptor in GaN, but the 
energy level is deep (Ea ==570 meV)[Pan97] and therefore should not be ionized at room 



88 

temperature. Based on the "NOG" scheme, the mechanism should be the same as the 
Ni/Au scheme discussed above, although the Zn might help increase the carrier 
concentration in the contact region at high temperatures. The main reason for improved 
contact performance with Ni/Zn-Au probably came from the limited time for native oxide 
to grow on the GaN and the use of high vacuum for metallization. Optimum contact 
resistance would not be predicted for this Ni/Zn-Au metallization because no reaction 
was found between Zn and GaN and no nitride forming component exist in the contact. 
A low resistivity (3.2x 10"^ Q-cm^) ohmic contact to p-GaN was produced with 
Ta/Ti metallization after a high temperature anneal (800°C for 20 min) [Suz99]. The 
authors postulated that Ta and Ti were able to remove hydrogen from Mg-H complexes 
and therefore reduced compensation of the acceptors. It was also found that a dual layer 
structure of both Ta and Ti formed better contacts than a single layer of either Ta or Ti, 
although both Ti and Ta were reported to have stronger binding energies with hydrogen 
than Mg. Although more hydrides are possible, it was reported in the literature that the 
common hydrides to Mg, Ti and Ta were MgH, TiH2 and TaHo.5, and the enthalpy for 
MgH (-0.77 eV/atom) was more negative than for TiH2 (-0.68 eV/atom) or TaHo.s (-0.417 
eV/atom) as shown in Table 5.1 [Fuk93]. This means that MgH is more energetically 
favored than the TiHa or TaHo.5. For these reaction products, Ta and Ti might not be able 
to reduce the amount of H in MgH. After a few days, the resistance of the Ta/Ti ohmic 
contacts increased to a much higher value. This was attributed by Suzuki et al to a 
reverse transport of compensating hydrogen from the Ti/Ta layers back to the interface 
region and recompensation of Mg acceptors. Contrary to the H mechanism, the "NOG" 
scheme would be consistent with a postulate that Ta and Ti would dissociate the GaN 



/ 89 

Table 5.1 Enthalpy and entropy of hydride formation [Fuk93] 



System 


AH"^P (eV/atom) 


AS"^'' (eV/atom) 


T(°C) 


Li -LiH 


-0.82 


-8.1 


600 - 900 


Na-NaH 


- 0.59 


-9.8 


500 - 900 


K-KH 


-0.61 


-10.1 


288-415 


Rb-RbH 


- 0.56 


- 10.2 


246 - 350 


Cs - CsH 


-0.59 


-10.2 


245 - 378 


Mg - MgH 


-0.77 


-8.1 


440 - 560 


Ca-CaH 


-0.95 


-8.4 


600 - 800 


Sr-SrH2 


-1.03 


-9.4 


<1000 


Ba - BaH2 


-0.91 


-8.6 


470 - 550 


Sc - ScH2 


-1.04 


-8.7 


>600 


Y-YH2 


-1.18 


- 8.7 


600 - 950 


YH2 - YH3 


-0.93 


-8.3 


250-350 


La - LaH2 


- 1.08 


-9.1 


600 - 800 


LaHz-LaHj 


-0.87 






Ce-CeH2 


-1.07 


-8.9 


600 - 800 


CeH2-CeH3 


-1.24 






Pr-PrH2 


-1.08 


-8.8 


600 - 800 


Nd-NdH2 


-1.10 


-8.8 


650 - 840 


Sm-SmH2 


-1.16 


-9.8 




Gd-GdH2 


- 1.02 


-7.9 


600 - 800 


Er-ErH2 


-1.18 


-9.4 




Ti (hep) - TiHa 


- 0.68 


-6 


<300 


Zr (hep) - ZrH2 


-0.98 


-9 


400 - 500 


Hf-HfH2 


-0.68 


-6 


600 - 900 


V-VH0.5 


-0.37 ■ 


-6.5 


0-100 


V-VH2 


-0.21 


-9 


50-120 


Nb-NbHo.65 


-0.48 


-8 


0-80 


Nb-NbH2 


-0.21 


-8 


25 


Ta-TaHo.5 


-0.41 


-6 


<50 


Mn(a)-MnH 


-0.11 


- 7 (assumed) 


450 - 730 


Ni-NiH 


-0.30 


- 7 (assumed) 


20 


Pd-PdHo.5 


-0.40 V 


-5 


-78-175 



■-:../.■ 



90 

and release N atoms. The released N atoms would increase the nitrogen chemical 
potential and result in reduced Vn concentrations, before formation of TaNx and TiNx 
compounds. /" , • ■ 

If the argument of hydrogen removal is correct, the sheet resistance of the p-GaN 
epilayer should decrease first and increase later during the contact degradation. Before 
contact degradtation, removal of hydrogen would reactivate the Mg acceptors, the carrier 
concentration in p-GaN would thus increased, and the sheet resistance should decrease. 
During the contact degradation, hydrogen move back into the p-GaN layer, compensating 
Mg acceptors and increasing the sheet resistance of the p-GaN. On the other hand, based 
on the "NOG" argument of nitrogen release and nitride formation, the sheet resistance 
would be predicted to decrease continuously because the released nitrogen would move 
to the contact interface and decrease the Vn concentration on the GaN surface. In 
practice, the sheet resistance was found to decrease continuously even after the 
degradation of the contacts [Mur99]. 

The differences in the thermodynamic and kinetic properties of Ta and Ti (Ta is 
stable and Ti is active ) might explain why Ta/Ti form better contacts than Ta or Ti 
individually. From the concept of "NOG", both metals might form good ohmic contacts 
in certain situations. In fact, Bour et al use Ti/Au as ohmic contacts to p-GaN in the as- 
deposited state for polycrystalline GaN LEDs [BouOO]. However, degradation of the 
contacts might proceed quickly. Low contact resistance are only observed if the 
measurement is done immediately after processing with the correct rapid thermal 
annealing (RTA) conditions. Observation of an increased resistance with time at room 
temperature can also be explained using the "NOG" postulate since formation of stable 



91 

nitrides would consume nitrogen atoms and create Vn, and thus increase the 
compensation of holes in the contact region. Some nitrides like TasNs are insulators, 
whose formation at the interface would increase contact resistance. The electrical 
conductivity of selected nitrides are listed in Table 5.2. 

For Ni/Mg/Ni/Si [You98] contacts, specific contact resistivities of «10'^ Q-cm^ were 
measured with a free hole concentration of 3x1 0'^ cm"^ These ohmic contacts also 
degraded after annealing at 500°C for 20min. The authors postulated a mechanism for 
contact formation which is known to occur for AuGeNi/GaAs contacts [Hol97] in which 
Ni first reacts to form a ternary compound with GaAs, and this ternary phase is 
subsequently dissociated by formation of binary NiGe, NixAsy and Au-Ga solutions. The 
release of Ga and As from the ternary phase allows solid phase epitaxial regrowth of 
GaAs in the presence of Ge dopant, leading to a n'^-GaAs regrowth layer and an ohmic 
contact [Hol97]. The authors in the study of Ni/Mg/Ni/Si postulated that regrowth of 
GaN and NiSi led to ohmic contacts. Using the principles of the "NOG" scheme, 
formation of an ohmic contact would result from dissociation of the GaN with Ni and 
formation of NiGax compounds and N atoms. This would increase the activity of N in the 
interfacial region, which might create a N-rich condition and a more highly doped p-type 
interfacial region. The formation of MgNx and SiNx would consume the extra N atoms 
and lead to the degradation of these contacts. 

The idea of interfacial reactions and control of vacancies, which is the basic tenet 
of the "NOG" scheme, can and does apply to ohmic contact to n-GaN as well as p-GaN. 
Lester, et al. [Les96] reported that aluminum produced an ohmic contact of 10" Q-cm to 
n-GaN. This is reasonable because of the matched work function of Al and GaN, plus Al 



4 



92 

Table 5.2 Electrical conductivity of selected nitrides [Sam75] 



Phase Resistivity, |aQ-cm T, °C Specific electrical Comment 

conductivity, Q"'-cm" 



CaaN 


5x10^ 


SraN 


2x10^ 


ScN 


25.4 


YN 


93 


LaN 


100 


CeN 


17 


PrN 


110 


NdN 


75 


SmN 


-120 


EuN 


-120 


GdN 


-200 


TbN 


-200 


DyN 


100 


HoN 


110 


ErN 


79 


TmN 


180 


YbN 


9x10^ 


LuN 


360 


ThN 




UN 


183 


NpN 


85 




-430 




380 


TiNo.79 


85 


TiNo.83 


78 


TiNo.87 


70 


TiNo.97 


40 


ZrNo.82 


30 


ZrNo.85 


28 


ZrNo.92 


20 





0.2 




5 


25 


39, 400 


25 


10, 750 


25 


10,000 


25 


58, 800 


25 


9,100 


25 


13,330 


25 


-8, 330 


25 


-8, 330 


25 


-5, 000 


25 


-5, 000 


25 


10,000 


25 


9,100 


25 


12, 650 


25 


5,550 


25 


111 


25 


2,780 


22 


5,460 


4K 


11,750 


82 K 


-2,380 


647 


2,630 


27 


11,750 


27 


12, 800 


27. 


14, 300 


27 


25, 000 


27 


33, 330 


27 


35, 700 


27 


50, 000 



Metal type 
Conductivity 



Curie Point 



93 



Table 5.2 - Continued 



ZrNo.97 


18 


HfN 


32 


V3N 


123.0±10 


VNo.71 


127 


VN0.80 


96 


VNo.85 


81 


VNo,87 


77 


VN0.93 


66 


VNo.96 


60 


Nb2N 


142 


NbNo.75 


109 


NbNo.80 


92 


NbNo.89 


72 


NbNo.95 


65 


NbN 


54 


TazN 


263 


TaN 


198 


TasNs 


~10'° 


CrjN 


84±5 


CrN 


640110 


M02N 


19.8 


NijN 


2.8x10^ 



27 


55, 500 


27 


31,300 


20 


8,140 


27 


7,880 


27 


10,400 


27 


12,350 


27 


13,000 


27 


15, 150 


27 


16, 660 


20 


7 042 


27 


9,180 


27 


10, 900 


27 


13,900 


27 


15,400 


27 


18,500 


20 


3,800 


20 


5,050 


20 


-10-^ 


20 


11,900 


20 


1,562 


20 


50, 500 


25 


357.1 



is also a nitride-forming metal which could create more nitrogen vacancies in the contact 
region. The contact resistance of Al/n-GaN increased by 50% upon annealing at 575°C. 
The postulated reason [Les96] was formation of a wide bandgap AIN layer at the 
interface. The specific contact resistance was improved to 8x10'^ Q-cm^ by the use of a 
Ti/Al bilayer metals annealed at 900°C, presumably due to the formation of TiN at the 
interface of Ti/Al (and Ti/Al/Ni/Au) contacts [Ruv96b, Fan96, Lin94]. Depletion of N in 
the GaN surface region with nitride-forming metals (e.g. Ti) would create more Vn, and 



94 

result in an n"^-GaN layer with improved electron tunneling probability and lower contact 
resistance. 

5.4 Experimental Studies 

Experimental data were collected in this study from Ni/Au, Ni/Ti/Au, Ag/Ti/Au, 
Ni/Al/Au, Pt/Au, Pt/Si/Pt/Au, Pt/Mg/Pt/Au contacts to p-GaN. Ni and Pt are used 
because they are gallide-forming metals, while Ag is a neutral metal with no reaction 
with GaN or the second Ti layer. Ti and Al are used because they are nitride-forming 
metals and also because the Al forms no hydride (so the possibility of hydrides might be 
excluded). Current-voltage (I-V) data showed that more current is obtained in the ternary 
or quaternary layer contacts designs using the "NOG" scheme, which is consistent with 
predictions. Although complete ohmic contacts were not formed with this metallization, 
the increased current shows the behavior expected from the "NOG" scheme. 

The p-GaN used in the Ni-based contacts were MBE epitaxy films, while those 
with Pt-based contacts were MOCVD GaN epilayers as described in Chapter 3. For Ni- 
based contacts, a Ni layer was deposited first in all contacts, and Ti and Al were used as 
nitride-forming metal. The thickness of Ni was varied from 20A, lOOA to 200A to study 
the effects of Ni thickness on the contact property and to find the best Ni thickness for 
later studies. The thickness of Ti or Al and of the protection Au is selected to be 500 A. 
The effects of annealing on the current through Ni/Au and Ni/Ti/Au contacts was studied 
at temperatures of 300°C, 500 °C or 800°C for times of 30 sec or 5 min in flowing ultra 
high purity N2 gas. 

The effects of Mg and Si layers on the contact performance were also studied with 
Pt as a gallide-forming metal because Si and Mg both are strong nitride-forming 



95 

elements. Mg is the p-type doping element in the p-GaN epilayer. Effects of annealing at 
temperatures of 600°C or 800°C were also studied for these Mg and Si contacts using 
flowing N2 as the protection gas. The p-GaN epilayers were grown with MOCVD 
method. 

5.4.1 Effects of Ti and Al as Nitride-Forming Metals 

Figure 5.2 shows the I- V data for as deposited Ni/Au, Ni/Ti/Au and Ni/Al/Au 

contacts. The film thickness are shown in the legend with the unit of A. Addition of Ti or 
Al to the Ni and Au contacts resulted in higher current levels by up to a factor of >4 at 
5V. Ti increased the current more than Al, although all contacts were still rectifying. 

Figure 5.3 showed the effects of Ni thickness on the contact I-V data in as- 
deposited state. The current through 20A Ni/500A Ti/500A Au was lower than that of 
500A Ni/500A Au contacts, while the current through the lOOA Ni/500A Ti/500A Au 
contacts was higher than that of the 500A Ni/500A Au contacts. The current for 200A 
Ni/500A Ti/500A Au contacts was similar to that from the 500A Ni/500A Au contacts. 

The effects of annealing on 500A Ni/500A Au and 200A Ni/500A Ti/500A Au 
contacts are shown in Figure 5.4. The as-deposited samples of both 500A Ni/500A Au 
and 2OOA Ni/500A Ti/500A Au had similar current levels. The current in 500A Ni/500A 
Au contacts increased continually upon annealing at 300 and 500"'C, while annealing at 
800°C decreased the current slightly. For 200A Ni/500A Ti/500A Au contacts, a 300°C, 
30 sec anneal resulted in slightly higher current, and a 5 min anneal at 300°C resulted in a 
four-fold increase in current at 5V. The contacts annealed at 300°C for 5min or 500°C for 
30sec exhibited similar currents, and current from both were comparable to or higher than 
that from a 500A Ni/500A Au contacts annealed at 500°C for 5min. After a 500°C, 5min 



96 





1 




0.8 




0.6 


< 

E 


0.4 
0.2 


■4-r 

c 

o 



-0.2 
-0.4 




-0.6 




-0.8 




-1 





A 








■W 


O500Ni/500Au 

n100Ni/1000Ti/500Au 

Al00Ni/1000AI/500Au 




., ..._... . 


1 1 I I 1 





-5-4-3-2-1012345 
Voltage, Volt 

Figure 5.2 1-V of Ni/Au, Ni/Ti/Au and Ni/Al/Au on p-GaN, as deposited state. 



annealing, the current through the 200A Ni/500A Ti/500A Au contacts decreased to a 
level which was close to that of as-deposited contacts. Annealing at 800°C for 5min 
resulted in a very low current (not shown). Similar degradation of contact versus 
annealing time and temperature was reported by Suzuki, et al [Suz99] and Ho, et al 
[Ho99a]. 

The higher current through as-deposited contacts, but serious degradation of the 
Ni/Ti/Au contact showed that while the nitride-forming metal (Ti) was helpful in 
reducing the contact resistance at the as-deposited states, it also led to the contact thermal 
instability. This would be consistent with a reaction between Ti and N and formation of 
TiNx which resulted in the generation of Vn's at the contact interfacial region. 



97 



< 



30 
20 
10 



1 



^^^T 



llk^':y 






^a 



-20 

-30 



t 



500ANi/500AAu 

D 20ANi/500ATl/500AAu 

A 100ANi/500ATl/500AAu 



-2 -1 1 2 

Voltage, Volt 

Figure 5.3 Effects of Ni thickness on the I-V of Ni/Ti/Au contact to p-GaN 



5.4.2 Effects of Si and Mg as Nitride-Forming Metals 

The contact thickness used in this part of study were 500A Pt/500 A Au, 1 OOA 
Pt/50A Si/ 500A Pt/500A Au and lOOA Pt/50A Mg/500A Pt/500A Au. An additional Pt 
layer was added to these contacts in hopes of achieving better protection of Si or Mg 
against the oxidation. 

The I-V data from as deposited 500A Pt/500A Au, lOOA Pt/50A Si/ 500A Pt/ 
500A Au and lOOA Pt/50A Mg/500A Pt/500A Au are shown in Figure 5.5. The lOOA 
Pt/50A Si/ 500A Pt/500A Au showed higher currents than Pt/Au contacts, while currents 



98 



< 
E 

♦J' 

c 
c 

k. 
k. 
3 

o 




-5 -4 -3-2-10 1 2 
Voltage, Volt 





0.8 




0.6 




0.4 


< 

E 


0.2 


c 

I 

3 

o 



-0.2 




-0.4 




-0.6 




-0.8 




o As deposited 
D 300C, 30sec 
A 500C, 30sec 
X300C, 5min 



-3-2-1012345 
Voltage, Volt 



Figure 5.4 Effects of thermal annealing on I-V data, (a) 500A Ni/500A Au and (b) 200A 
Ni/500A Ti/500A Au contact. 



99 





10 T 

8- 






._ J 


< 

E 


6 

4- 

2 






•^ 


c 


0- 




^^jtfsmss 


53^^^ 


3 

o 


-2- 

-4- 
< 


aO |-|J 


p>^ 










oPt/Au 




■«-^ 






D Pt/Si/Pt/Au 




-8- 


J 






A Pt/Mg/Pt/Au 










-10 i 1 1 

-5 -4 -3 


— 1 — 1 1 
-2 -1 


12 3 4 5 








Voltage, 


Volt 





Figure 5.5 I-V curves of 500A Pt/500A Au, lOOA Pt/50A Si/500A Pt/500A Au 

idlOOA Pt/50A Mg/500A Pt/500A Au contact on p-GaN, as-deposited 



anc 




OPt/Aj 

DR/a/R/Pu 

APt/lV^Pt/Au 



-5 4 -3 



-2-10 1 2 
Volt^e,Volt 



Figure 5.6 I-V curves of 500A Pt/500A Au, lOOA Pt/50A Si/500A Pt/500A Au 
and lOOA Pt/50A Mg/500A Pt/500A Au contact on p-GaN, 600°C for 
Imin annealing 



■■*• t i ' 



T\^**^ 



/•» >i-H 



4- -•':4 V 



100 



40 





30 




20 


< 


10 


c 
£ 

3 

o 



-10 




-20 




-30 




-40 




tUaAAAAA AAAAAAAAAAAAAA 



yi^AAAAAA A AAAAAAAAAAAAAAA 



Pt/Au 

Pt/Si/Pt/Au 

Pt/Mg/Pt/Au 



-5-4-3-2-1012345 

Voltage, Volt 

Figure 5.7 I-V curves of 500A Pt/500A Au, lOOA Pt/50A Si/500A Pt/500A Au 
and lOOA Pt/50A Mg/500A Pt/500A Au contact on p-GaN, 800°C for 
Imin annealing 



from lOOA Pt/50A Mg/500A Pt/500A Au contacts were lower than those in 500A 
Pt/500A Au contacts. All contacts were rectifying. 

The change in current levels versus voltage data after high temperature annealing 
at 600°C or 800°C for 1 min. are shown in Figures 5.6 and 5.7, respectively. The I-V data 
for the 500A Pt/500A Au and lOOA Pt/50A Si/500A Pt/500A Au became more linear 
while the current for lOOA Pt/50A Mg/500A Pt/500A Au became very low. Annealing at 
800°C decreased the current level in the lOOA Pt/50A Si/500A Pt/500A Au and lOOA 
Pt/50A Mg/500A Pt/500A Au contacts, while that from lOOA Pt/500A Au contact 
remained relatively, but lower in magnitude than as-deposited contacts (Figure 5.5). 



101 

AES depth profiles for 500A Pt/500A Au contacts are shown in Figure 5.8. 
Compared to the as-deposited state (a), diffusion of Pt and Au was found after the 
annealing. Figure 5.9 showed the depth profiles of lOOA Pt/50A Si/500A Pt/500A Au 
contacts. Compared to profiles of lOOA Pt/500A Au contacts, less interfacial diffusion 
was found after amiealing at 600T (comparing Fig. 5.9b with 5.8b), but the Si peak 
height decreased compared to as-deposited data. Less diffusion of Pt was expected due to 
formation of Pt silicides [Oka90]. Annealing further at 800°C promoted more interfacial 
reactions and diffusion so that Si accumulated at the contact surface. Oxygen also 
accumulated on the sample surface, alsong with higher N levels than those for the as- 

deposited state. 

For lOOA Pt/50A Mg/500A Pt/500A Au contacts, as shown in Figure 5.10, the as- 
deposited state exhibited a Mg peak at the contact/p-GaN interface, along with higher 
level of oxygen. Annealing at 600°C leveled the Mg peak at the interfacial region, but a 
huge accumulation of Mg and O was found at the contact surface. The inset in Figure 
5.10-b shows the Mg and O peaks. The N level was higher than that of lOOA Pt/50A 
Si/500A Pt/500A Au contacts, and both were higher than for the 500A Pt/500A Au 
contacts. Annealing at 800°C made the Mg and accumulation wider at the contact 
surface, and the interfacial region was wider than for 500A Pt/500A Au and lOOA Pt/50A 
Si/500A Pt/500A Au contacts. 

From the above results, Mg is not good as the nitride-forming metal because of thermal 
instability. The Mg contacts became almost insulating after annealing at 600°C and 
800°C for Imin. This is understandable because Mg metal reacts strongly with oxygen. 






102 



12000 



9000 



•f 6000 



(a) 

Au 



R 






Ga 




200 



400 600 

Sputter Time, sec 



t fOAi -*- -^ ■ 

800 1000 



1200 



12000 



n 



9000 



f 6000 

X 

S 3000 
Q. 



(b) 







Ga 



/^' 



200 



400 600 800 

Sputter Time, sec 



1000 1200 



12000 




200 400 600 800 1000 1200 

Sputter Time, sec 



Figure 5 8 AES depth profile of Pt/Au contact on MOCVD p-GaN. (a) 
As-deposited; (b) 600°C for Imin; (c) 800°C for Imin 



103 



12000 



2 9000 - 



•ff 6000 



0) 




200 400 600 800 1000 1200 

Sputter Time, sec 



12000 




200 400 600 800 1000 1200 

Sputter Time, sec 



12000 




200 400 600 800 1000 1200 

Sputter Time, sec 



Figure 5.9 AES depth profile of Pt/Si/Pt/Au contact on MOCVD p-GaN. (a) As- 
deposited; (b) 600°C for Imin; (c) 800°C for Imin 



■ * -* * 



: "'f / 



104 



12000 



9000 N'V^/. / '^''^ 




200 400 600 800 

Sputter Time, sec 



1000 1200 



12000 



ir.'^^\->\/M,. 



Pt •40000 I 






SI 100 2O0 




200 400 600 800 

Sputter Time, sec 



1000 



1200 



12000 




400 600 800 

Sputter Time, sec 



1200 



Figure 5.10 AES depth profile of Pt/Mg/Pt/Au contact on MOCVD p-GaN. (a) As- 
deposited; (b) 600°C for Imin; (c) 800°C for Imin 



105 



1.2 



0.8 



< 


0.4 


E 




♦J 




c 





g 




3 




o 


-0.4 




-0.8 




-1.2 




-5 -4 -3 



O 500Ni/500Au 
n100Ni/1000Ti/1000Au 
A 100Ag/1000Ti/1000Au 



-2-1012 

Voltage, Volt 



Figure 5.1 1 I-V of 500A Ni/500A Au, lOOA Ni/IOOOA Ti/IOOOA Au and lOOA 
Ag/IOOOA Ti/IOOOA Au on p-GaN, as deposited state 



5.4.3 Neutral Metals 

The effects of a neutral metal, Ag, upon I-V data were compared with those from 
Ni as the first contact layer. The Ag did not react with either GaN or Ti metal [Oka90], so 
its role in the Ag/Ti/Au contact is only to separate the Ti from GaN. 

The I-V data (Figure 5.11) showed that the current through lOOA Ag/IOOOA Ti/ 
lOOOA Au contacts was below that through lOOA Ni/IOOOA Ti/IOOOA Au contacts, but 
higher than the 500 A Ni/IOOOA Au contact. The I-V plots were highly non-linear due, 
presumably, to the absence of interfacial reactions between Ag and GaN to stabilize the 
Ga atoms and release the N atoms, in contrast to the reaction between Ni and GaN. 






106 




♦ In contact 

D Ni/Ti/Au contact 



-2-10 1 2 
Voltage, Volt 



Figure 5.12 Comparison of In metal and lOOA Ni/500A Ti/500A Au as ohmic 
contact for Hall measurement 



Another interesting comparison was between lOOA Ni/500A Ti/500A Au and the 
plain In contacts (Figure 5.12). Because good ohmic contacts of high current and linear I- 
V plots to p-GaN are not routinely available, most researchers use plain In and the Van 
der Pauw configuration for Hall measurements. The lOOA Ni/500A Ti/500A Au contacts 
were prepared with the shadow mask described in Chapter 3, while the In contacts were 
soldered to p-GaN surface. The In (with a diameter of ~1 mm) contact had a larger 
contact area than the lOOA Ni/500A Ti/500A Au contacts (with a diameter of 0.4 mm). 
The lOOA Ni/500A Ti/500A Au contacts exhibited higher current and more linear I-V 
data, even though their area was smaller. 



107 

5.5 Discussion 
The "NOG" contact scheme proposed in this section was shown to be 
quaHtatively consistent with the behavior of metal contacts on GaN. The reactions 
between the contact metals schemes and GaN could determine the electrical properties of 
the contacts. Presumably, the reactions between gallide-forming metals and GaN gettered 
excess Ga in the interface region, thereby decreasing the nitrogen vacancies and 
increasing the hole concentrations. Reactions between nitride-forming metals and GaN 
increased the nitrogen vacancy concentrations and enhanced the electron concentrations 
in the contact region. 

Consistent with the principles of "NOG", the first Ni layer in both Ni/Ti/Au and 
Ni/Al/Au contacts was a gallide forming metal which could release N atoms upon 
reaction with GaN. This Ni layer also act as a barrier by slowing down the rate of nitride 
compound formation, reducing the extent of reactions between N or Al and Ti. This 
would maintain a higher activity of N at the GaN surface to reduce the Vn concentration, 
and result in less acceptor compensation. The observed higher current is consistent with 
this qualitative prediction. A high contact resistance for thin gallide-forming metal (20A 
Ni) could result from the extent of interfacial reaction being too limited to affect Vn 
concentration or the Ni layer is too thin to prevent quick formation of stable nitride 
(TiNx) and thus to increased Vn concentrations. 

The driving force of metal reactions to Ga and GaN can be calculated using the 
equation of 

AG = AH-T-AS (5.1), 





108 






Table 5.3 Calculated driving forces of metal reactions to Ga and GaN 




Reaction 


AG300 (kJ/mol) 


AGiooo (kJ/mol) 
-2.5 


TcforAG = 0(K) 
752.7 


Ga+'/2N2 = GaN 


-77.6 


Co + Ga = CoGa 


-39.4 


-35.6 


7592.6 


Ni + Ga = NiGa 


-36.1 


-34.1 


12758.6 


Pt + Ga = PtGa 


-40.1 


36.7 


666.5 


Co + GaN = CoGa + 'A N2 


38.2 


-32.9 


676.5 


Co + GaN = CoGa + N 


510.9 


439.9 


5331.9 


Ni + GaN = NiGa +I/2N2 


41.5 


-31.4 


698.7 


Ni + GaN = NiGa + N 


514.2 


441.4 


5242.2 


3 Ni + GaN = NisN + Ga 


104.8 


91.6 


5831.1 


Pt + GaN = PtGa +I/2N2 


5.5 


-67.4 


352.7 


Pt + GaN = PtGa + N 


477.4 


402.5 


4763.4 


Cr + GaN = CrN + Ga 


-19.1 


-32.3 


<0 


2Cr + GaN = CrN + CrGa 


-22.3 


-5.7 


1237.9 


3 Cr + GaN = Cr2N + CrGa 


87.1 


153.5 


<0 


2 Cr + GaN = Cr2N + Ga 


47.7 


-15.3 


830.5 


Ta + GaN = TaN + Ga 


-137.3 


-124.6 


7905.8 


Ti +GaN = TiN + Ga 


-223.3 


-215.0 


18987.4 


3Mg + 2GaN = Mg3N2 + 2Ga 


-211.2 


-346.6 


<0 


3Si + 4GaN = Si3N4 + 4Ga 









where AG is Gibbs free energy change, ^1// and AS are enthalpy and entropy change for 
the reaction, and Tis the temperature in Kelvin [Deh88]. Using the data provided in the 
literature [Boe88, Deh93], the driving forces for reactions between typical metals and 
GaN are calculated as in Table 5.3. The negative values showed that reactions between 
gallide-forming metals of Co, Ni and Pt with Ga and GaN are thermodynamical favorable 
if the produced nitrogen exists in the form of nitrogen gas (N2), or combines with nitnde- 



' ■ ■ 109 

forming metals (Cr, Ta and Ti, etc.) to form nitrides, but is not favorable if the released 
nitrogen exists in atomic state (N). 

The "NOG" contact scheme uses these properties of gallide formation and uses a 
nitride-forming metal to keep the nitrogen activity high in the contact region. However, 
thermal stability became a serious problem in the "NOG" contact metals tested in this 
study. Thermal instability resulted from the large diffusion distances of nitride-forming 
metals, as illustrated below. 

The temperature dependence of the diffusion coefficient in a solid is well 
described by a semi-empirical formula [Deh96] 

\ D = D,-cxp{-£^) (5-2) 

where Do is a diffusion factor, Q is the activation energy, R is the universal gas constant 
and ris temperature in Kelvin [Deh93]. The characteristic diffusion distance, A, is 
usually written as: 



X = ylA-D-t (5-3) 

where the D is given above in equation (5.2). 

The diffusion data of a few metals in nickel are Hsted in Table 5.4 [Gri97]. The 
diffusion coefficients at 973 K and the characteristic diffision distance for / = 1 hour were 
also calculated and are shown in the table. At 973 K (700°C), the characteristic diffusion 
distance is much larger than the Ni contact thickness used in this work. The diffusion data 
of N is not available. Using the diffusion data of O (due to similarity in atomic radii and 
atomic affinity between O and N), the diffusion distance is estimated to be between 18.5 
A and 0.355^m for a time period of 60 min as shown in Chapter 4.The large diffusion 
distances of contact components are probably a major cause of the thermal stability 



no 



Table 5.4 Calculation of diffusion characteristic distance of selected metals in nickle 
[Gri97] 



Metal 


Temperature 
Range, K 


Do, cm^'/sec 


Q, kJ/mol 


D at 973 K 


>., nm 


Al 


1070-1250 


1.1 


249 


2.6x10-'^ 


193.23 


Au 


1200-1400 


2 


272 


2.4x10-'^ 


59.06 


Cr 


940-1170 


0.03 


171 


1.3x10-" 


4329 


Mg 


1070-1250 


2.3x10-^ 


131 


1.5x10"'^ 


1497 


Si 


1079-1250 


10.6 


271 


1.5x10-"* 


149.42 


Ti 


1400-1600 


11.1 


322 


2.0x10-'^ 


1.716 



problem in the "NOG" scheme. To alleviate the problem, better metal layer should be 
found to prevent serious diffusion of these nitride-forming metals. 

5.6 Summary 
New principles to be used to select metal layers used for contacts to p-GaN were 
proposed and called Nitride-forming metal Over Gallium-forming metals — "NOG". The 
"NOG" contact scheme is based on reactions between the contact metals and GaN. 
Gallide-forming metals can bond to Ga and release N atoms, while addition of nitride- 
forming metals should result in increased N concentrations and activity in the contact 
interface before the formation of stable nitrides. These N atoms could fill nitrogen 
vacancies, with a consequent higher hole concentration and better ohmic contact. These 
results were qualitatively obtained in comparison of as-deposited Ni/Ti/Au, Ni/Al/Au, 
Pt/Si/Pt/Au and Ag/Ti/Au contacts with Ni/Au and Pt/Au contacts. 

Explanations based on "NOG" principles of published data as well as data fi-om 
the current study were discussed. In the "NOG" scheme, metals were classified as 



Ill 

gallide-forming, nitride-forming or neutral metals. Published results for ohmic contacts to 
p- or n-GaN were compared to "NOG" principles. Higher currents through contacts with 
both gallide- and nitride- forming metals were presented (e.g. Ni/Al/Au and Ni/Ti/Au), as 
compared to Ni/Au contacts with only a gallide-forming and neutral metal. Mg and Si 
were also studied as nitride-forming elements. The Si contacts showed increased current, 
but Mg decreased the current levels versus Pt/Au contacts. These results were explained 
based on metallurgical reactions between the contact elements and GaN. 

Even though the use of both gallide and nitride forming metals slightly reduced 
the contact resistance, poor thermal stability was observed for metallization based upon 
the "NOG" scheme. The annealing temperatures degraded the "NOG" contacts more 
quickly than those having no nitride-forming metals. This was explained based on the 
formation of nitrides which would decrease the activity of N atoms in the contact region. 
This would create more nitrogen vacancies and further reduce the already low hole 
concentrations. 



CHAPTER 6 
EFFECTS OF Ni CAP LAYER ON THIN Ni/Au CONTACTS TO p-GaN 



6.1 Introduction 

Recently, a new approach to reduce the contact resistivity for p-GaN was reported 
to result from annealing Ni/Au contacts in an oxygen ambient (pure oxygen or air). Both 
the sheet resistance and specific contact resistance were found to be reduced [Koi99]. A 
low pc of 4x10'^ Q-cm^ was obtained with 50A Ni/50A Au contacts after annealing in air 
[Ho99b]. While this value would be sufficiently low to be used in patterned small area 
contacts, most contact resistances were much higher (« 10"^ Q-cm ). The mechanisms to 
explain these low contact resistivities were postulated to be reduction of acceptor- 
compensating hydrogen [Koi99], or formation of aNiO layer [Ho99b, Che99, HolOO]. 
NiO is a p-type wide bandgap semiconductor. It can potentially reduce the Schottky 
barrier height {<f)b) at the p-GaN/contact interface if it has the expected large work 
function (or Eg + %). In the study by Ho et al [Ho99b], it was found that equal thicknesses 
of Ni and Au resulted in the best electrical properties. Low values of specific contact 
resistance of 4 x 10'^ Q-cm^ and 8 x 10"^ Q-cm^ were obtained from contacts of 50A 
Ni/50A Au and 200A Ni/200A Au. The contact resistance increased for thicker contacts. 

Because of its wide bandgap (4.0 eV) [Huf92], NiO also is transparent to visible 
light allowing it to be used as transparent electrodes for optoelectronic devices and 
window coatings [Sat93]. Since pc « 10"* Q-cm^, which is too high to be used for small 



112 



113 

area patterned ohmic contacts, a transparent contact is used in LEDs over the entire 
emitting surface. With such a large area, the contribution of the contact to the total device 
resistance is very small. The contact contribution to total device resistance increases for 
thin contacts and decreases for thick contacts. But the light transmittance decreases as the 
contact thickness increases according to the equation [Hum98]: 

1 = 1^-6-'"' (6.1) 

where h is initial light intensity, /is transmitted light intensity, a is the light attentuation 
constant and d is distance the light travels through a material. Therefore the thickness of 
the tranparent contact is a compromise between these opposed trends for resistance 
versus transmittance. 

Another consideration for the thickness of the transparent Ni/Au contacts is the 
surface roughness. The reported root mean square (RMS) surface roughness of GaN 
epilayers grown under typical MOCVD processing conditions is normally alOnm 
[CaoOO]. As shown in Figure 4.1 l-(a), with a RMS roughness of 2.7nm, the peak-to- 
valley distance is around 18.8nm for MBE-GaN. This is greater than the thickness of 50A 
Ni/50A Au contacts. In a recent report on polycrystalline GaN LEDs, the surface 
roughness was much larger; and contact coverage was one of the major factors limiting 
device performance [BouOO]. To make this thin Ni/Au contact scheme more robust, the 
thickness needs to be increased. 

In this work, an outer Ni cap-layer was added to Ni/Au contacts making them 
Ni/Au/Ni contacts. The effects of the additional Ni layer on contact resistance and light 
transmittance were studied. Annealing of this contact in O2 transformed the extra Ni into 
a transparent NiO layer, with little negative effects on the light transmittance or light 



114 

transmittance. Thus the Ni cap-layer can increase the contact thickness, achieve better 
coverage, and still keep the same light transmittance and contact resistance. Optimization 
of the Ni/Au ratio was studied by changing the Au layer thickness in the sandwich 
Ni/Au/Ni structures. 

The thickness of the Ni layer was kept constant at 50A for comparison with 
published work. Two gold thicknesses of 50 A and lOOA were used, so the contacts 
studied were 50A Ni/50A Au, 50A Ni/50A Au/50A Ni, 50A Ni/IOOA Au/50A Ni 
(referred to as 50/50, 50/50/50 and 50/100/50 hereafter). Ho et al [Ho99b] reported that 
the best Ni/Au contact resistance was obtained by annealing at a temperature of 500°C for 
lOmin. After an anneal of 600°C, the contact degraded significantly. Due to the increased 
Ni thickness in this work, the best anneal temperature and time for Ni/Au/Ni were 
expected to be higher and longer, so anneal temperatures of 500, 550 and 600°C and 
times of 1, 5; 10 and 30min were selected. Specific contact resistance and light 
transmittance were measured, and SEM, AES and XPS were used to study 
microstructural evolution and interfacial reactions. 

6.2 Contact Electrical Properties 
The electrical properties of the contacts were characterized by the current levels at 
a constant voltage, and by specific contact resistance. The specific contact resistance for 
these contacts was calculated using the circular transfer length method (CTLM) patterns 
shown and discussed in Chapter 3. Annealing time was first kept constant at lOmin to 
study the effects of annealing temperatures. After selection of the best annealing 
temperature, the annealing time was optimized. 



115 




-1 -0.5 0.5 1 

Voltage, Volt 

Figure 6.1 Effects of anneal temperature on I-V of 50/50 contact 



Formation of NiO was important to the contact properties, and the electrical 
conductivity of NiO was found to be related to its composition [Sat93]. To determine 
whether the annealing temperature and time influenced the compositon and/or electrical 
conductivity of the contact pad, the resistance within a single patterned contact (not 
between two patterned contacts) was measured with a two point probes approach. The 
resistance calculated from these data have contributions from pNio» Pau and poaN- 



6.2.1 Annealing Temperature 

Figure 6.1 depicts the effects of annealing temperature on the I-V curves for a 
constant time of lOmin for 50/50 contacts. The as deposited contact was rectifying. After 
armealing at 500°C for lOmin, the I-V curve became more but was not completely linear. 
After armealing at 550°C, the contact became less resistive with straight, linear I-V with 



116 

higher current levels. After 600°C annealing, the I-V data were still linear although the 
current level decreased. 

The I-V data obtained from the 50/50/50 contacts are similar to data from 50/50 
contacts, except that the current level is slightly higher for the as deposited state. As 
compared to the as-deposited sample, both the current and linearity increased as the 
annealing temperature was increased from 500°C to 550°C to eOO^C. Completely linear I- 
V data were obtained at 600°C. However, a current of only ~1.5mA at IV was obtained, 
which is lower than the current in 50/50 contacts. The current at 550°C was between that 
for samples annealed at 500°C and 600°C. Annealing at 700°C, as compared to 600°C 
sample, significantly decreased the current levels. 

For 50/100/50 contacts, increased annealing temperatures increased the current, 
with the maximum current obtained after 600°C anneal (0.4mA at IV). However, 
completely linear I-V curves were not obtained for this scheme, and the current through 
as-deposited contacts was lower as compared to that from 50/50 and 50/50/50 contacts. 

The values of pc are shown in Figure 6.2 for 50/50, 50/50/50 and 50/100/50 
contacts versus anneal temperature. For non-linear I-V curves, the resistance at 0.2 V was 
used for this calculation. The lowest contact resistivity for the 50/50 contacts was 5.3x10' 
'*Q-cm^ at 550°C, and the resistivity was 1.52x10"^ at 600°C. For the 50/50/50 scheme, 
the lowest specific contact resistance was 6.2x10"* Q-cm^ at 600°C. Annealing at 700°C 
increased the specific contact resistance by more than one order of magnitude. The 
50/100/50 contacts all had high resistivity (>10'^ Q-cm^) at all annealing temperatures. 
From these data, an annealing temperature of 600°C was selected for 50/50/50 and 
50/100/10 contacts. 



117 



10 



0.1 



E 
o 

G 

o 
u 

c 
re 

m 

w 
o 
DC 



5 0.01 

c 

o 

o 

o 

1 0.001 

0) 

Q. 
(0 



0.0001 



400 



-X- 50/50 
-O— 50/50/50 
-D- 50/1 00/50 




500 600 700 

Anneal Temperature, °C 



800 



Figure 6.2 Effects of annealing temperature on the specific contact resistance of 
the 50/50, 50/50/50 and 50/100/50 contacts 



6.2.2 Effects of Annealing Time 

For the 50/50 contacts, annealing at 600°C for 1 min. increased the current 
slightly, and completely linear I-V data were obtained with lOmin anneals. Increasing 
armealing time to >30min resulted in lower, nonlinear current. 

I-V data for 50/50/50 contacts showed similar trends as found for 50/50 contacts 
except for a smaller decrease in current levels for 30min anneals. 

The trends were similar for 50/100/50 contacts, except that a decrease in current 
was not found for annealing at >30min. The current continued to increase for all times at 



118 

T = 600°C for 50/100/50 contacts, although its values were always lower than those for 
the 50/50 or 50/50/50 contacts. 

The specific contact resistance results versus time at 600°C are shown in Figure 
6.3. The 50/50 and 50/50/50 contacts show minimum specific resistance of 1.52 x 10" Q- 
cm^ and 6.16 X 10"* Q-cm^ respectively, after 10 min. anneals. The specific resistance of 
50/100/50 contacts continually decrease with increased annealing time, with a minimum 
value of 1.6 X 10'^ Q-cm^ after anneahng for 30min. 

The effects of annealing time at 600°C on contact pad resistance is shown in 
Figure 6.4. The resistance for the 50/50 contact pads increased sharply for lOmin anneal, 
but remained constant between annealing times of 10 and 30min. The 50/50/50 resistance 
of contact pads increased continuously. The contact pad resistance remained at a low 
value for the 50/100/50 contact, although it increased slightly with annealing time. 

6.2.3 Effects of O7 Flow Rate 

Effects of oxygen flow rates (60 to 130 liter/min) on the 50/50 and 50/50/50 

contacts were studied after 10 min anneals at 500°C. A MANOSTAT® gas flow meter 
was used to set flow rates of 60 liter/min or 130 liter/min. In 50/50 contacts, the lower O2 
flow rate resulted in smaller currents, while the larger O2 flow rate enhanced the current 
level between two contact pads. Increased current between pads with larger O2 flow rates 
was also found for 50/50/50 contacts. 

For the 50/50/50 contact schemes. The contact pad resistance increased at higher 
O2 flow rates through the furnace, contrary to the decrease of specific contact resisitvity 



119 



10 



C3 

oT 
u 

c 

s 

V) 

« 

0) 

u 
B 

c 
o 
o 



o 

o 0.001 

0) 

a. 



0.1 



0.01 



0.0001 



-X- 50/50 
-0— 50/50/50 
HD- 50/1 00/50 




10 20 

Anneal Time, min 



30 



J 

40 



Figure 6.3 Effects of annealing time on the specific contact resistance of 50/50, 
50/50/50 and 50/100/50 schemes annealed at 600°C. 



observed with an increased O2 flow rates. The reasons for change in specific contact 
resistance and contact pad resistance are unknown. 



6.3 Light Transmittance 
Applications of GaN in optoelectronic devices like light emitting diodes (LEDs) 
and laser diodes (LDs) are important. Because the specific contact resistance of current 
ohmic contacts to p-GaN is too high (-10"^ to 10"^ Q-cm^), a large contact area is used to 
reduce the contribution of the contact to the total resistance. Light emission through the 



120 



500 



400 



S 300 
o 

c 

•2 200 

(A 
O 

100 



50/50 

50/50/5 

50/100/5 




10 20 

Time, min 



30 40 



Figure 6.4 Effects of anneal time on resistance of contact pads at 600°C 

p-contact is necessary, requiring a transparent contacts. To maximize the emission, the p- 
type contacts must be as transparent as possible. The results of annealing temperature, 
time and change of O2 flow rates on light transmittance is reported below. 



6.3.1 Effects of Annealing Temperature 

Figure 6.5 shows the change of light transmittance in the 50/50, 50/50/50 and the 
50/100/50 contacts with annealing temperatures. In the as deposited state, the 50/50 was 
the thinnest contact, so the transmittance was higher than that of the 50/50/50 and 
50/100/50 contacts. Similarly, the 50/50/50 scheme had a better transparency than the 
50/100/50 scheme in the as-deposited state, and this situation remained true after 
annealing at 500°C for lOmin. For all three contacts, an increase of annealing temperature 
improved the light transmittance, with the maximum values obtained at 600°C. 



121 



100 



80 



b 




c 






60 


g 




tf) 




^ 


40 






H 




4i* 




Jl 




o> 




_l 


2U 



■ 50/50 
- 50/50/50 
■50/100/50 




1 1 1 1 1 1 I i~ 

100 200 300 400 500 600 700 

Temperature, "C 

Figure 6.5 Comparison of light transmittance at X = 450 nm. Increased annealing 
temperature led to higher light transmittance (armeal time = 10 min). 

Surprisingly, 50/50/50 contacts showed the highest light transmittance, which was more 
than 93%. hi contrast, 87% and 75% were measured for 50/50 and 50/100/50 contacts, 
respectively, after annealing at 600°C for lOmin. 

6.3.2 Effects of Annealing Time 

The relation between light transmittance and armealing time at 600°C is shown in 
Figure 6.6. The light transmittance reached a maximum at lOmin for the 50/50 (88%) and 
50/50/50 contacts (94.5%), and decreased slightly as the annealing time increased to 
30min. The 50/100/50 contact reached a value of light transmittance of 76% at lOmin and 
continued to increase to 79% for 30min. At 30 min , the 50/50/50 contact (88%) was still 



more transparent than the 50/50 contact (83%). 



122 



100 



80 



s 




e 




s 


60 


§ 




c 




SS 


40 


1- 




4-1 




£ 




O) 




J 


20 












50/50 

50/50/50 

50/100/50 



10 20 30 

Anneal Time, min 



40 



Figure 6.6 Effects of annealing time on light transmittance at A, = 450 nm 



6.3.3 Effects of O? Flow Rates 

The effects of O2 flow rates on the light transmittance are shown in Figure 6.7 for 
anneal temperature of 500°C. The annealed contacts show higher transparency than as- 
deposited contacts (47% for the 50/50 contacts and 36% for the 50/50/50 contacts). The 
high O2 flow rate (130 liter/min) results in higher light transmittance in both contact 
schemes. The 50/50 (50% at low flow rate and 63% at high flow rate) contacts were more 
transparent than the 50/50/50 contacts (46% at low flow rate and 55% at high flow rate) 
at both flow rates, which is contrary to the relative transmittance of 50/50 versus 
50/50/50 contacts annealed at 600°C (see Figure 6.6). 



123 




As 

deposited 



low O2 
rate 



High O2 
rate 



Figure 6.7 Effects of O2 flow rates on light transmittance for 500°C, lOmin 
anneals (low rate = 60 liter/min, high rate =130 liter/min). 



6.4 Microstructure Characterization 
6.4.1 SEM 

Figure 6.8 to 6. 1 1 shows the microstructures imaged with backscattering electrons 
for the 50/50, 50/50/50 and 50/100/50 contacts on p-GaN. All as-deposited and short time 
annealed (600°C for 1 min.) films are featureless to SEM. Annealing times of 5min 
results in the formation of a light colored network structure, and an increase to lOmin 
enhanced this network structure. Energy dispersive spectroscopy (EDS) point chemical 
analysis (Figures 6.12 and 6.13) shows the light region is Au rich while the dark region is 
Ni rich. The feature sizes are < 0.5 fim. Because the excitation volume diameter (~1 |im) 
is larger than the average feature size, the presence of Au in EDS spectra fi-om the dark 
regions (Figure 6.13) is not unusual. 



_^'-- ^j»i 



124 








(b: 




^^^^O^vl^ 




^ * 


1 




* 




mm^ 




* 






1 






\ 
'i 




-^ 




■l*,'l' 



,r >..;■£ 
















Figure 6.8 Microstructure of the 50/50 contact after annealing. SEM 

backscattering electron image after (a) 600°C for 1 min; (b) 600°C 
for 5 min; (c) eOO^C for 10 min and (d) 600°C for 30 min. 



125 




Figure 6.9 Microstructure of the 50/50/50 contact after annealing. SEM 

backscattering electron images after (a) 600°C for 1 min; (b) 600°C for 5 
min; (c) 600°C for 10 min and (d) 600°C for 30 min. 



126 




Figure 6.10. SEM backscattering electron image of same sample as in 

Figure 6.9-(d) but at a higher magnification showing the Au film 
is still continuous. 



As the annealing time is increased from 1 min to 5 min, a very uniform microstructure 
developed in the 50/50 contact scheme. The 50/50/50 contact exhibited a similar 
microstructure except smaller pores were observed in the white Au matrix. In the 
50/100/50 contact, the microstructure consists of larger pores and better defined 
boundaries between white and dark regions. These microstructures were maintained in all 
three contacts during increased anneal time from 5 to 10 min, except in all instances there 



127 










> 



Figure 6.1 1 Microstructure of the 50/100/50 contact (SEM backscattering image) 
after annealing at (a) eOO^C for Imin, (b) 600°C for 5min, (c)600°C for 
lOmin, and (d)600°C for 30min. 



128 






cps 



80- 



60- 



40- 



20- 



<■! 




EnetjyCkeV) 



Figure 6.12 EDS analysis of the light region in Figure 6.16-c showing a larger 
Au/Ni peak ratio. 




EnergyO<e\/) 



Figure 6.13 EDS analysis of the dark region in Figure 6.16-c showing a larger 
Ni/Au ratio. 



cps 

250- 
200- 
150- 
100- 



50- 



03 



Au 



I u 




129 



Nl 



■I I I"" 



Au 



Ga 



■V^^ 



Au Au 



10 



Energy (keV) 
Figure 6.14 EDS analysis of spherical particle in Figure 6.16-d showing a very 
large Au/Ni ratio. 



appear to be some isolated growth of Ni-rich dark pores to larger sizes (~0.5^m). As the 
annealing time increased from 10 min to 30min, minor changes were observed in the 
50/50/50 and 50/100/50 samples, but numerous bright sub-micron spherical particles, 
surrounded by large dark areas, precipated on the 50/50 samples (Figure 6.8d). Point EDS 
analysis of these particles showed that they were Au and Ga rich, with no N signal 
detected (Figure 6.14). The excess Ga must result from dissociation of GaN. This result 
suggests that the thin Ni/Au contacts are not stable after 600°C, 30min armeal. 

For the purpose of understanding the microstructural evolution, single layers of Ni 
or Au are deposited on GaN or sapphire substrates and annealed in oxygen ambients. The 
results show that Au without the presence of Ni forms submicron islands on both GaN 
and sapphire (Figure 6.15), similar to the Hght precipitates in Figure 6.8-(d). Annealed Ni 
with no Au cap layer forms a featureless, continuous NiO film (Figure 6.16). 



130 





Figure 6.15 Annealed Au film on (a) MOCVD-GaN; (b) sapphire (600°C, 
lOmin). 



131 




Figure 6.16 Annealed (600°C, lOmin) Ni film on GaN showing no 

formation of islands in contrast to Fig.6.15. The white particle is 
contamination used for focusing of the SEM. 



6.4.2 AES Survey and Depth Profiling 

Figure 6.17 shows AES surveys fi"om the 50/50 contact for different processing 
conditions, and Figures 6.18 and 6.19 shows similar surveys for 50/50/50 and 50/100/50 
contacts. The Au signal was only found at the surface of as-deposited state, but was not 
found at the surface after annealing at 500°C or 600''C for lOmin or 600°C for 30min. For 
the 50/50/50 contact scheme, higher signal of oxygen was found on the as-deposited 
surface than for 50/50 scheme due to the presence of original Ni, and the oxygen signal 
increased further as the annealing temperature and time increased. The 50/100/50 scheme 
showed similar trends to the 50/50/50 contacts. 

The AES depth profile for the 50/50 scheme is shovra in Figure 6.20, while those 
fi-om 50/50/50 and 50/100/50 contacts are shown in Figures 6.21 and 6.22, respectively. 



132 




2500 



400 800 1200 1600 2000 
Kinetic Energy, eV 




400 800 1200 1600 2000 
Kinetic Energy, eV 



12500 




2500 



400 800 1200 1600 2000 
Kinetic Energy, eV 




400 800 1200 1600 2000 
Kinetic Energy, eV 



Figure 6.17 AES surface spectra from 50/50 contacts, (a) As-deposited; (b) 500°C 
for lOmin; (c) 600°C for lOmin; (d) 600°C for 30min. 



For as deposited 50/50 contacts, the interface region between Ni and GaN was relatively 
sharp and the oxygen level was very low. Annealing at 600°C for lOmin changes the 
surface to be only NiO with no Au detected. Compared to the as-deposited state, the 
interface between NiO/Au/GaN became very diffuse, either from diffusion or the 
microstructure shown in Figure 6.8. This is consistent with published reports [Ho99] that 
Au stayed in the NiO/GaN interfacial region. The oxygen level increased and followed 



133 



■r -<*» 




400 800 1200 1600 2000 
Kinetic Energy, eV 




400 800 1200 1600 2000 
Kinetic Energy, eV 



2500 




2500 



400 800 1200 1600 2000 
Kinetic Energy, eV 




400 800 1200 1600 2000 
Kinetic Energy, eV 



Figure 6.18 AES surface spectra from 50/50/50 contacts, (a) As deposited; (b) 
500°C for 1 Omin; (c) 600°C for 1 Omin; (d) 600°C for 30min. 



the Ni signal except for lower oxygen concentrations in the interfacial region. Increased 
annealing temperatures (up to 600°C) increased the 0/Ni ratio in the top layer. Increased 
annealing time to 30 min resulted in no change in the distribution of these elements. 
There may have been more Ga was also found in the contact interfacial region with 
higher annealing temperatures and times. A Ga plateau in the contact interface at all 
temperature and time suggests formation of a Ni-Au-Ga compound (Figure 6.20-(b), (c) 



134 



2500 




2500 



-2500 



400 800 1200 1600 2000 
Kinetic Energy, eV 




400 800 1200 1600 2000 
Kinetic Energy, eV 



2500 




400 800 1200 1600 2000 
Kinetic Energy, eV 



Figure 6.19 AES surface spectra from 50/100/50 contacts, (a) As- 
deposited; (b) 600°C for lOmin; (c) 600°C for 30min. 



and (d)). In the 50/50/50 and 50/100/50 schemes, Figure 6.21 and 6.22, the Ni and Au 
signals showed the two Ni layers and the sandwiched Au layer in the as-deposited states. 
Annealing at 600°C for lOmin leveled the Ni humps and made the Au move into the 
interfacial region. A large oxygen to nickel ratio was found as the annealing temperature 
or time increased. The increase in the annealing time to 30min at 600°C did not result in 
significant changes in the depth profiles. 



135 



40000 



3 

i 30000 






20000 



a. 



10000 




90 180 270 360 450 
Sputter Time, sec 



40000 



3 

i 30000 



■« 

z 

a 
o> 
a. 



20000 



10000 




90 180 270 360 450 
Sputter Time, sec 



40000 



3 

«, 30000 

I 20000 

10000 



(c) 






■\Ni 




Ga 


y^uvV 


X"^^^ 


N 


^.tif-^^^ 





90 180 270 360 450 
Sputter Time, sec 



40000 




90 180 270 360 450 

Sputter Time, sec 



Figure 6.20 AES depth profile of 50/50 contacts, (a) As-deposited; (b) 500°C for 
lOmin; (c) 600°C for lOmin anneal; (d) 600°C for 30min. 



136 



40000 



3 

i 30000 



'S 

X 

(S 

0) 

0. 



20000 



10000 




90 180 270 360 450 
Sputter Time, sec 



40000 




90 180 270 360 450 

Sputter Time, sec 



40000 



3 

«. 30000 

a 

I 20000 

n 
« 
a. 

10000 



(c) 


Ni 








K 




y 


^y^~^ 


c 

' ■ ■■■ 


y. 


ys 

^ 




N 



90 180 270 360 450 

Sputter Time, sec 



40000 



3 

«. 30000 

** 

I 20000 

ni 
O 

a. 

10000 



W) 








)^ 


N": 




Ga 




^; 


^ 


/^^^^^ 


- 




Au vj^ 


N 


C 


y^ 


"•^"^'^ 





90 180 270 360 450 
Sputter Time, sec 



Figure 6.21 AES depth profile of 50/50/50 contact, (a) As deposited; (b) 500°C 
for lOmin; (c) 600°C for lOmin anneal; (d) 600°C for 30min. 



6.4.3 XPS Analysis 

To show the evolution of chemical states associated with the contact Ni metal and 

interfacial reaction products, XPS surveys were taken from the 50/50 and 50/50/50 

contacts after annealing at 600°C for lOmin. The bonding data were collected from the 

fresh surface after sputtering for a short time with an Ar ion beam (see Chapter 3). 



137 



50000 



3 
O) 


40000 
30000 


o 

r 


20000 


Q. 


10000 




90 180 270 360 

Sputter Time, sec 



450 





50000 






3 




n 


40000 


^4-* 






30000 


0) 




z 


20000 


Ji 






10000 



" (b) 


NI 


Ga 


N 





80 



180 270 360 

Sputter Time, sec 



450 



_ 50000 
n. 40000 
1, 30000 
X 20000 

S 10000 


' (c) 

c 


''^ 


Ni Ga 


D. 



< 




> 


80 


180 270 360 450 
Sputter Time, sec 



Figure 6.22 AES depth profile from 50/100/50 contacts, (a) As deposited; (b) 
600°C for lOmin; (c) 600°C for 30min. 



138 

Figures 6.23 and 6.24 shows the Ni2p peaks at different sputtering times in 50/50 
and 50/50/50 contacts. For 50/50 contacts, the as received surface showed the Ni2p 
binding energy of 855eV. After 5min sputtering, this energy decreased to 853.5 eV. After 
15min sputtering, this energy decreased further to 852.5eV, and only a small amount of 
Ni was detected after 15min sputtering. The presence of NiO is supported with the 
presence of shake up structures and chemical shift in Figure 6.23-(a) and (b) [Mou95], 
The lack of shake up peaks and energy shift indicates that the XPS spectrum in Figure 
6.23-(c) is from metallic Ni. For 50/50/50 contact surfaces, the binding energy of Ni2p 
was again high and the shake-up stucture showed the compound to be NiO. The binding 
energy decreased after sputtering for 5min, and corresponds to the binding energy of 
NiO, and then decreased further to 853eV representing a mixture of oxide and metallic 
bond at the contact/p-GaN interface. 

Changes in binding energy was also observed for the 01s signal (Figure 6.25 and 
6.26) and Ga2p (Figure 6.27 and 6.28) peaks. The high energy shoulder on the 01s peak 
from both 50/50 and 50/50/50 contacts suggest the presence of hydroxides, consistent 
with higher binding energies on as-annealed surfaces from the Ga2p and 3d spectra. 
Sputtering for 5min, removed the high binding energy 01s shoulder and reduced the 
Ga2p and 3d binding energies, consistent with removal of a surface layer of Ni(0H)2. 
The lower binding energies for Ni 2p and 3d after sputtering to near the contact/GaN 
interface and the loss of shake-up structure in the Ni 2p spectra showed the presence of a 
Ni-Au-Ga metallic layer at the interface, consistent with the AES depth profile plateaus 
obvious in Figures 6-20, 6-21 and 6-22. 



.* < ft 



1 

100000 
90000 




139 




-(a) 




ili2p"^ 


yj 80000 


- 


f 


m 70000 

^ 60000 


"~~^ 


Shake-up peaks 


\ 


50000 
40000 


1 


1 1 -I — 


, 



UJ 

uT 



100000 



80000 



60000 



40000 



tu 
2 



100000 



80000 



60000 



40000 



900 890 880 870 860 850 840 
Binding Energy, eV 




900 890 880 870 860 850 840 
Binding Energy, eV 




900 890 880 870 860 850 840 
Binding Energy, eV 



Figure 6.23 XPS spectra of Ni2p from 50/50 contact after annealing at 600°C for 
lOmin. (a) As-annealed surface; (b) 5 min sputtering; (c) 15min sputtering. 



140 



90000 

80000 

^ 70000 

uT 

Z 60000 

50000 

40000 




900 890 



880 870 860 
Binding Energy, eV 



850 840 




900 890 



880 870 860 
Binding Energy, eV 



90000 



850 840 




900 890 880 870 860 850 840 
Binding Energy, eV 



Figure 6.24 XPS spectra of Ni2p from 50/50/50 contact after annealing at 600°C 
for lOmin. (a) As annealed surface; (b) 5 min sputtering; (c) 15min 
sputtering. 



141 



40000 



S 



30000 



20000 



10000 




540 



535 530 525 

Binding Biergy, eV 



520 



40000 



30000 



ui 



20000 



10000 




540 



535 530 525 

Binding Energy, eV 



520 




30000 



540 



535 530 525 

Binding Biergy, eV 



520 



Figure 6.25 XPS spectra of 01s from 50/50 contact after annealing at 600°C for 
lOmin. (a) As-received Contact; (b)5 min sputtering; (c) 15min sputtering 



142 



30000 



20000 



UJ 

uT 



10000 




540 



535 530 525 

Binding Energy, eV 



520 




20000 



540 



535 530 525 

Binding Energy, eV 



520 



30000 



^ 20000 



10000 




540 



535 530 525 

Binding Energy, eV 



520 



Figure 6.26 XPS spectra of 01s from 50/50/50 contact after annealing at 
600°C for lOmin. (a) As-received contact; (b)5 min sputtering; (c) 
1 5 min sputtering 



143 



UJ 

uT 



100000 



90000 



80000 



70000 




1130 1125 1120 1115 

Binding Energy, eV 



1110 




ui 

S" 90000 



85000 



80000 



1130 1125 1120 1115 

Binding Energy, eV 



1110 




80000 



1130 1125 1120 1115 

Binding Energy, eV 



1110 



Figure 6.27 XPS spectra of Ga2p from 50/50 contact after annealing at 

600°C for lOmin. (a) As received contact; (b) 5 min sputtering; (c) 
15min sputtering 



144 




80000 



1130 1125 1120 1115 

Binding Energy, eV 



1110 




80000 



1130 1125 1120 1115 
Binding Energy, eV 



1110 



180000 

160000 

uj 140000 

^ 120000 

100000 



80000 




1130 1125 1120 1115 1110 
Binding Energy, eV 



Figure 6.28 XPS spectra of Ga2p from 50/50/50 contact after annealing at 600°C for 
lOmin. (a) As-received Contact; (b) 5 min sputtering; (c) ISmin sputtering 



145 

6.5 Discussion 

The mechanism by which oxidized thin Ni/Au films lead to ohmic contact to p- 
GaN has been discussed [Ho99b]. The effects of thick versus thin p-NiO were 
considered. With a thick NiO layer, a large portion of the applied voltage is dropped 
across the p-NiO. The impedance increases with increasing reverse bias voltage because 
of the barrier and the resistive NiO layer. With a thin NiO layer, the existence of a hole 
notch close to p-GaN makes the carrier transport easily in the contact region due to a field 
emission mechanism, thus formation of an ohmic contact is possible. However, the 
microstructure development and the effects of a Au-Ni-Ga interfacial layer has not been 
included in the model. 

In Ho's model [Ho99b], the distribution of barrier between GaN and p-NiO was 
calculated. With a carrier concentration of 2 x lO'^ cm"'', the Fermi level in p-GaN was 
calculated to be 0.13 eV above the top of the valence band at 300 K. For undoped p-NiO, 
the carrier concentration was estimated to be 1 x lo'^ cm''', and the Fermi level was 
calculted to be 0.5 eV above the top of the valence band. From these parameters, and by 
equaling the built-in potential across the p-NiO/p-GaN isotype heteroj unction to the work 
function difference between p-GaN and p-NiO (before the formation of a junction), a 
barrier value of 2.47 eV was obtained. The distribution of this barrier was further 
calculated to be 2.415 eV in p-NiO and 0.055 eV in p-GaN. The structure of this contact 
to p-GaN was regarded as Au/p-NiO/p-GaN as shown in Figure 6-29. An ideal ohmic 
contact would be formed between Au and p-NiO because the work function of Au ((J)au = 
5.10 eV) is larger than that of p-NiO ((|)p.Nio = 4.9 eV). A hole notch is formed at the 
NiO/p-GaN hetereoj unction because of the large band offset between p-NiO and p-GaN. 



146 



NiO p-GaN 



^f zz. 




VACUUM 
LEVEL 



Figure 6.29 Energy diagram of oxidized thin Ni/Au contact to p-GaN 



Because the Au network dispersed in the matrix of NiO is covered with NiO, the energy 
diagram presented in [Ho99b] for thin p-NiO contacts to p-GaN is modified as in Figure 
6.29 by the addition of the NiO levels to the left of Au. The above discussion successfully 
explains the mechanism of ohmic contact between thin oxidized Ni/Au and p-GaN, 
however, no explanation was given for ignoring Fermi level pinning in NiO or GaN. 
Fermi level pinning of GaN was described in Chapter 2 [Ren98]. 

The validity of ignoring the pinning of the Fermi level in these contacts might be 
justified by the reaction between Ni and GaN. These reactions to form Ni-Ga compounds 
(e.g. Ga4Ni3 and Ga3Ni2 [Che99]) can lead to intimate interface bonding between Ni and 
GaN. This would reduce the density of surface states from dangling bonds, which is 
believed to be the major source of Fermi level pinning [Bar47]. 

Besides the reaction between Ni and GaN, the match of lattice constants is also 
important to reduced contact resistivity by reducing the defects in the contact region. The 
lattice constants of GaN, Ni, NiO and Au are listed in Table 6.1 [Edg98, Liu89]. The 
lattice constants of NiO and Au differ by only -2.5%, and a reaction between Ni and 



147 



Table 6.1 Lattice constants (A) of components in oxidized Ni/Au contacts to p-GaN 



GaN Ni NiO Au 

a = 3.189 3.499 4.177 4.065 
c = 5.185 



GaN to form a compound might also improve the -12% mismatch of the lattice constants 
between GaN and Ni. The large mismatch in lattice constants between Ni and NiO can be 
modified continuously by change of the oxygen composition. In this way, a lattice 
matched contact system of NiO-Au-NiO-Ni-GaN could be formed. Which would be 
expected to reduce the interfacial defect. In fact, orientation relationships are found 
among NiO, Au and GaN [Che99], which is: 

NiO(l 1 1)//Au (1 li)//GaN(0002) and 
NiO[liO]//Au [liO]//GaN[1120]. 
Lack of development of strong texture or even epitaxy at the contact interface could 
explain why sputter deposited thick NiO/Au resulted in poor contact properties [MaeOO]. 

The AES data (Figures 6.17 to 6.22) showed the as-deposited surface were metals 
(either Au or Ni) with a NiO layer formed at all the contact surfaces after annealing. The 
Au rearranged to the contact interfacial region, and most of the Ni diffused out to form 
NiO. This is consistent with published work [Che99]. After addition of the Ni cap layer, 
the inner Ni layer still difftised out to the surface region to react with oxygen, but the 
contact properties changed, and the mechanism need to be discussed. 

The Ni cap-layer was converted to NiO after annealing in O2. As discussed above, 
NiO is a wide-bandgap semiconductor which is semi-insulating, so the conductivity of 
the contact pad (contact sheet resistance) would decrease, as observed in this study. 
However, the specific contact resistances of annealed 50/50/50 contacts were close to 



148 



those of the 50/50 scheme for the optimum amieahng condition (Figure 6.2). From 
calculation of the depletion width, W, in the tunneling current regime [SzeSl] 



w= \^'^^- 



qNs ' 



(6.2) 



we have the relation of 



n,= 



1-e, 



(6.3) 



where W is the total depletion distance, Cs is semiconductor permittivity, Vti is built-in 
potential, q is magnitude of electronic charge, and Nb is either acceptor or donor impurity 
density. This shows that an increased NiO thickness (equal to an increase of the depletion 
distance W) should form a larger barrier. For light Transmittance, the Au layer thickness 





Yl2 

/ 


Phase 1 


Ql^- 


X 


^ 


Phase 2 


Yl3 \ 


4' - 


) 


Q2 


Y23 


Phases 



I , Figure 6.30. Schematic diagram of interface equilibrium between 
three phases 



149 

was the same for both contact schemes, but the increased NiO thickness should normally 

be expected to decrease the amount of transmitted light 

The Ni/Au ratio of unity [Ho99b] was reported to yield the lowest contact 

resistivity. With the sandwich Ni/Au/Ni contacts, the critical parameter appears to be the 

Au thickness rather than the ratio of Au/Ni thickness. 

The microstructure evolution of the Ni/Au and Ni/Au/Ni contacts may be 
explained using the concept of surface tension. Advanced modeling of thin film 
morphologies is based on work by Srolovitz and Safran [Sro86] which shows the Ni cap- 
layer is effective in improving the thermal stability of the contacts. Last, reaction kinetics 
are considered to show that the microstructural morphology differences are possible 
reasons for different light transmittance in these samples. 

The surface tension in thin films is shown in Figure 6.30. A three phase 
equilibrium situation is considered. The simple Dupre equation [Mur75] governing the 
balance of interfacial free energies for three phases of 1, 2 and 3 in equilibrium in the 
absence of interfacial torques is: 



Yii Ym Yn 



(6.4) 



sinQ, sinQj sinQj 
where yu, yu and yzs represent the phase boundary free energies between phases 1 and 2, 
1 and 3, and 2 and 3, respectively. The Di. ^2 and /?? are the dihedral contact angles 
measured in phases 1, 2 and 3, respectively. Phase 1, 2 and 3 may be solid, liquid or 
vapor, or any variation thereof 

To show the effects of Ni cap-layer on the microstructure evolution, the surface 
tensions values involved are listed in Table 6.2. No surface tension data were found for 
Ni/GaN, Ni/sapphire, Au/GaN and Au/sapphire interfaces. These values were calculated 



150 
Table 6.2 Values of surface tension ofNi and Au [Mur75b, Cha87] 



System 


Surface Tension (JW) 


AuA'^ac 


1.658 


Au/Ni 


0.114 


Au/NiO 


-2.0 


Au/Au (grain boundary) 


0.438 


NiA^ac 


2.533 


Ni/NiO 


0.392 


Ni/Ni (grain boundary) 


0.958 


Ni/Ni (twin boundary) 


0.052 


NiO/NiO 


1.00 



at 400°C, which is similar to the temperature used in this work. The surface tension 
decreased after replacing the gold/vacuum interface, AuA^ac, with a gold/nickel, Au/Ni, 
interface. 

In Ni/Au contacts with no Ni cap-layer, the three phases of Au, air and Ni exist in 
the contact. With the Ni cap layer, these three phases become Ni, Au and Ni. The contact 
angle of Au on Ni in these two situations was calculated using the values listed in Table 
6.2. Before annealing, the equilibrium contact angles of the Au on Ni (73 in Figure 6.30 
where Ni is phase 3 and Au is phase 2) were calculated to be 0°, for the contact with and 
without a Ni cap layer. This explains the results shown in Figure 6.16-19 which showed 
that Au and Ni spreaded on the whole contact surface and the as-deposited surfaces were 
featureless. After annealing in O2, the surface layer became NiO for both situations and 
the Au moved to the interfacial region. Because a value of surface tension for GaN is 
unknown, the Au was assumed to be sandwiched between two NiO layers as reported by 
[Ho99b]. Using the above equation, this contact angle (Q3) was calculated to be 150°, 



151 

which means the Au would not remain flat, but islands or porous networks should form 
as shown schematically in Figure 6.30. 

This is consistent with the formation of islands of Au in the annealed single Au 
film on GaN and sapphire, as seen in Figure 6.15. However, this argument does not 
explain the formation of cavities (holes), rather than "islands" in the annealed Ni/Au or 
Ni/Au/Ni contacts. • 

Hole formation in thin films have been widely reported. Chao et al [Cho87] 
observed porosities in thin Ni/Au metallization layers annealed in oxidizing environments 
from 250 to 400°C. Kane et al [Kan66] observed hole formation in gold films deposited 
onto or between cadmium sulphide (CdS) layers upon heating to 450°C in the 
atomosphere of helium. Gimpl et al [Gim64] reported agglomeration in gold and nickel, 
and formation of holes was also reported by Jaeger et al [Jae69] in silver films prepared 
under ultrahigh vacuum conditions followed by exposure to air. 

Numerous mechanisms have also been suggested for the formation of porosity in 
metallization, such as grain boundary and bulk interdiffusion [Bac80, NakSO], generafion 
of stresses [Hum71], and grain boundary grooving [HumSl]. Grain boundary grooving 
driven by capillarity (surface tension) was shown to be detrimental to the stability of thin 
gold films. The effects of grain boundary grooving on the film continuity was modeled 
by considering the surface and grain boundary energies and conserving the total volume 
of metal [Sro86]. 

The stability of the pores in thin film contacts was evaluated using the Srolovitz 
and Safran model [Sro86]. For pores with a radius rp existing in a film of thickness a and 



?1 -'i- 



152 



AE 




a=10nni,p= 457.4 



Figure 6.31 AE versus p (= x/K) calculated for 50/50 and 50/50/50 
contacts, respectively. See text for assumption. 



an average distance between pores of2R, the energy change between a porous and a 
continuous thin film, AE, can be written as: 



AE 



-4 



2 , ^4 



where 



n-a -y^ 3 + 4-lnp-4-p +p 



y R^ 



-P.P- 



(6.5), 



R 



and /v is the surface energy of gold in ambient, /s = 7sv- 7/s, Ysv is the substrate-ambient 
surface energy, and y/s is the film-substrate interfacial energy. 



1^3 

Because all the samples are heat treated in oxidizing ambients, a NiO overlayer 
exists on the Au film as shown with AES data. The samples were actually Au films 
sandwiched between Ni at the substrate and NiO at the surface, just as was true in the 
work of Chao et al [Cha87]. The Srolovitz and Safi-an model can be extended to the 
present thin film contacts by replacing the Au/ambient surface energy by the Au/NiO 
solid surface energy, i.e., by replacing the ambient vapor by a solid NiO layer, i.e. the 
substrate/vapor energy ysv is the interfacial energy of Ni and NiO. To calculate p, the 
values of /au/nio, Yau/ni, and Ynunio in Table 6.2 are used. The results in [Ho99b] showed 
that for 50A Ni/50A Au metallization, the contact thickness is around 330A after 
oxidation, with no explanation given for the expansion. However it seems reasonable to 
assume this results fi-om conversion of Ni to NiO, reactions between Ni-GaN-Au, and 
changed microstructure. Assuming the same ratio for the 50A Ni/50A Au contacts also 
applies to Ni/Au/Ni contacts, the thickness is calculated to be 66OA. For comparison, the 
initial thickness of 150 A and 100 A are also included for calculation. Substituting these 
surface energies into the expression for p and taking 2R = 0.2 [im (as measured in Figure 
6.8), a - 330A and 100 A for the 50A Ni/50A Au and a = 660 A and 150 A for the 50A 
Ni/50A Au/50A Ni, the value of J3 is calculated to be 42, 457.4, 10.5 and 203.3, 
respectively. 

The plots of zl^'vs r/R (p) are shown in Figure 6.30. Fot/3~ 10.5, the energy of 
the system increases as the pore radius increases, so spontaneous growth of pores is not 
expected in contacts with Ni cap layers. For /?- 42, 203.3 and 457.4, the energy of the 
system is reduced as pores grow, and it reaches a minimum at r/R = 0.55. Spontiiueous 
growth of pores is expected since growth will lower the system fi-ee energy. For 



154 

Ni/Au/Ni, smaller pore size is expected as compared to Ni/Au contacts with no Ni cap 
layer. 

Contrary to the data in Figure 6.31 for p = 10.5, pores do grow in contacts with Ni 
cap layers, as shown in Figure 6.9 and 6.10. This inaccurate prediction probably results 
from oversimplification of the actual situations. The Srolovitz and Safran model assumes 
a single value for grain size and idealized geometry. While these facts do not invalidate 
the model, they do cause inaccuracies. Stress in the thin films may also influence the 
growth of pores, but the model does not consider strain energy which would certainly 
modify the curves in Figure 6.31. 

It should also be recognized that the Srolovitz and Safran model assumes that no 
defects exist in the gold film prior to heat treatments. Gold is well known to contain 
pinholes and pores in the as-deposited films which might grow to bigger sizes as 
observed in this study. However, the porosity in this work is not believed to come from 
inherent defects in the deposited Au film because the pore density is extremely high and 
uniform over the sample. Random, large inherent defects were observed, but they did not 
lead to such uniform microstructures as shown in Figure 6.8 to 6.1 1. 

To explain the increased light transmittance in 50/50/50 contacts, the 
microstructural morphology must be considered. With the thin Au film, the out-diffusion 
of Ni from the inside layer still allows the Ni-Au-Ni structure to exist for at least a short 
time. This results in a contact angle (A in Figure 6.29) of 153.64°(using the twin 
boundary energy in Table 6.2) is equilibrium is approached. Because of limited transport 
distances, the density of nucleation sites required to approach this equilibrium structure 
would be large and would result in the very uniform microstructure as observed. Once 



NiO 



155 



.«-f f 



/"^Au 


\^ 




Au ^ 


\f^ 


\( 


Au ( 


"V „ 




GaN 







NiO 



(a) 



fS, 








A 


/' 


Au 




GaN 



NiO 



(b) 



\^ 










(^ 








c\ 


/ 


Au 


\ 


GaN 



(c) 



Figure 6.32 Schematic diagram of contact microstructure showing (a) 50ANi/ 50A Au, 
(b) 50A Ni/50A Au/50A Ni and (c) 50A Ni/IOOA Au/50A Ni contacts. In 
50A Ni/50A Au contacts, the Au network pore area is less, causing more 
light absorption in the thin Au regions. In 50A Ni/50A Au/50A Ni contacts, 
the area of pores is larger with a subsequent thicker Au network, resulting in 
less absorption, more transmittance of light. Finally in 50A Ni/IOOA 
Au/50A Ni contacts, a thicker Au network and lower pore area 
microstructure leads to less transmitted light. 



" c 156 ■ 

this microstructure had developed after an anneal of 600°C forlOmin, it did not 
significantly change with longer time annealing (to 30min). The lack of large Au islands 
as shown in Figure 6.8-(d) shows much better thermal stability for 50/50/50 contacts 
versus 50/50 contacts. 

In the SEM pictures (Figure 6.8 to 6. 1 1 ), the light color region corresponds to Au 
and the dark region is NiO. In the 50/50 contact, the Au film stayed as a continuous layer 
with dark "pores" (the term used in [Cha87]). The Au-rich regions occupied a larger 
fraction (-70%) of surface area as compared to the pores (-30%, Figure 6.8-(c)). In the 
50/50/50 contact, however, it was hard to tell which component was dominant between 
the Au network and "pores" (-50% for each component. Figure 6.9-(c)). The Au and NiO 
interlace with each other and form a very uniform mixture, but the Au was still 
continuous. The Au network covered a smaller fractional area of the surface as compared 
to the 50/50 contact. Since the original Au thickness was the same in both contact 
schemes before annealing, the Au areas in the 50/50/50 should be thicker than those in 
the 50/50 contacts after armealing. 

The schematic diagram of the microstructure is shown in Figure 6.32 for the three 
contact schemes. Although the NiO would be thicker in the 50/50/50 scheme, the 
increased Au thickness would still allow significant current through the contact because 
of its continuous network. As a result, similar contact specific resistance was obtained for 
50/50/50 contacts as in 50/50 contacts. On the other hand, the Au film in the 50/100/50 
contact was thick and occupied a large area again (Figure 6.1 1 -(c)), and the 
microstructure was similar to that of the 50/50 contact (Figure 6.8-(c)). 



157 

The schematic diagram of microstructures in Figure 6.32 explains the higher light 
transmittance in the 50/50/50 contacts. If light transmittance through the Au film is 
ignored (the characteristic penetration depth is 150A in Au film at a wavelength of 589.3 
nm [Hum92]), then the light only goes through the "dark" pores (as seen in the SEM 
pictures). The annealed 50/50/50 contacts cause a larger area percent of surface to be 
pores and would allow more transmittance of light. In practice, some light still goes 
through the thin Au film although it experiences attenuation (to 37% of the initial 
intensity if the Au film is equal to the characteristic thickness of 150A). In the 50/100/50 
contact, the increased Au feature size was due to the thicker Au film resulting in lower 
transmittance. 

In this argument, light scattering from the microstructure has been ignored. Since 
the features are typically less than half the wavelength of light, this seems reasonable. 
However, diffraction effects can still change the intensity and direction of light scattering 
from such small patterns. This needs to be explored in fiiture research. 

6.6 Summary 
Formation of low resistance "transparent" ohmic contacts using oxidized thin 
Ni/Au films on p-GaN was a major achievement. Using this scheme, the specific contact 
resistance was reported in the literature to be as low as 4x 10"^ Qcm^, although the best 
results in this study was 10"* Q-cm^. While thinner ohmic contact would be more 
transparent allowing better light output fi-om high brightness LEDs or LDs, thin contacts 
may exhibit thermal instability, coverage problems, and higher contact sheet resistances. 
To increase the contact thickness while retaining good electrical properties and high light 



• >r 158 

transmittance, a thin Ni capping layer was placed on top of Au/Ni/GaN contacts.The Ni 
cap-layer was shown to transform into transparent NiO upon annealing and result in 
higher light transmittance (-93% versus 85% for 50/50 contacts) while retaining the same 
low contact resistance (-lO""* Q-cm^). Thicker Au with the same 50ANi cap-layer resulted 
in higher contact resistance (-10"^ Q-cm^) and lower light transmittance (60%). Based on 
SEM backscattering images, the microstructures consist of a continuous Au network with 
Ni-rich pores reaching the GaN surface. The area fraction of pores was larger for 
50/50/50 contacts versus 50/50 contacts, which explains the higher light transmittance. 
Using AES and XPS depth profiling, the Ni layer was found mainly to diffuse out and 
form NiO on the surface, but a fraction of the Ni metal reacted with GaN. 

The development of the microstructure consisting of a Au network with Ni-rich 
pores was modeled based on the concept of surface tension and grain boundary grooving. 
The Srolovitz and Safran model predicted that without a Ni cap-layer, pores would lower 
the system free energy and grow automatically to a limiting critical radius. With the 
addition of a Ni cap-layer, the pores increased the system free energy and should not 
grow. However, it was speculated that strain energy still allowed them to grow, but to a 
smaller, more uniform size. This is consistent with the observed improved thermal 
stability of Ni capped contacts. 






CHAPTER 7 
CONCLUSIONS 



Formation and improvement of ohmic contacts to p-GaN has been studied using I- 
V data from patterned TLM contacts. H2O2 was used in passivating the surface of p-type 
MBE GaN, resulting in higher conductivities. A "NOG" scheme was proposed to help 
understand formation of ohmic contacts to p-GaN, and addition of a Ni cap-layer to thin 
transparent Ni/Au contact was shown to increase the light but still maintain a low contact 
resistance. • 

Using the H2O2 treatment, an increase in the carrier concentration up to 1 00% for 
MBE grown p-GaN doped 2 x lO^' cm"^ with Mg. Peroxide: water solutions ranging from 
1:5 to pure (37%) H2O2 were tested for times ranging from 30sec to 60min. For treatment 
time, SOsec < t < SOmin, increased carrier concentrations were demonstrated for 1 : 1 
solutions. For t = 60min, reduced carrier concentrations were found. H2O2 treatment did 
not affect n-GaN or p-GaN grown by MOCVD. Possible explanations for the effects of 
H2O2 were discussed, with the best model being reduction of acceptor compensation by 
either hydrogen removal or reduction of nitrogen vacancies. Nanopipe transport of H202 
species are thought to help the kinetics. 

The "NOG" contact scheme was developed based on the knowledge of 
metallurgical reactions between transition metals and GaN. The approach of using a 
nitride-forming metal over a gallide-forming metal can modify the thermodynamic 
activity of N and Ga near the interface. This in turn can change the near-surface point 

159 



-; 160 

defect concentrations, particularly the N vacancies and excess Ga. The principles of this 
contact scheme were shown to qualitatively agree with literature data from Ni/Au, Ni/Zn- 
Au, Ta/Ti, and Ni/Mg/Ni/Si contact schemes. In the current experimental study, the 
"NOG" scheme was used to design metallizations of Ni/Ti/Au, Ni/Al/Au, Pt/Si/Pt/Au and 
Pt/Mg/Pt/Au. The addition of Ti, Al and Si as nitride-forming metals to the Ni or Pt 
gallide-forming metal led to improved current transport across the interfacial barrier. Ti 
was shown to be better than Al as the nitride-forming metal in as-deposited contacts. 
Compared to Ni/Au, four times more current was measured in Ni/Ti/Au contacts to p- 
GaN after annealing at 300°C for 5 min. However, selection of metallization schemes 
based on the "NOG" model was not able to decrease the contact resistivity to a level low 
enough for device applications. This was ascribed to interfacial diffusion which 
consumes the extra N atoms in the interface region. Besides, addition of Ti metal led to 
lower thermal stability of the Ni/Au contact. Addition of Mg to Pt/Au contacts also led to 
serious thermal stability problems. 

Addition of a Ni cap-layers to thin transparent Ni/Au ohmic contacts to p-GaN 
was aimed at improving the contact microstructure. This Ni cap layer was oxidized to a 
transparent NiO layer and resulted in a higher light transmission (-93%), same low 
contact resistivity (-10"^ Q-cm^) and better thermal stability. 

The changes in the electrical and optical properties as well as the microstructure 
evolution after adding the Ni cap-layer were explained using surface tensions and the 
total system energy of thin film structures. With a capping layer of Ni, the increased 
contact angles from interface tension led to a finer structure and allowed more light 



161 



transmittance through the contact. Because of the increased thickness, high anneal times 
and temperatures were needed to reach the optimized properties. 



•:^' \ ' -.yi 



CHAPTER 8 
FUTURE WORK 



GaN and related materials have presented exciting opportunities in materials 
research and device engineering. This has become an example of the beneficial interplay 
between science and technology. The continued advancement of new devices made from 
GaN requires better understanding of the mechanisms controlling the formation of ohmic 
contacts to p-GaN, and this understanding needs to result in lower ohmic contact 
resistance. 

For the H2O2 solution treatment of p-GaN, the most important thing should be to 
find a good procedure to clean the newly formed oxides after the H2O2 treatment, with 
the purpose of removing the excess Ga atoms but not increasing the oxides at the contact 
interfacial region. More work in finding better parameters will be definitely helpfiil in 
improving of the contact properties. 

For the "NOG" scheme, work must be continued to find a contact scheme that 
provides both low resistance and good thermal stability to p-GaN. Titanium has been 
shown to form contacts with higher currents than the Ni/Au scheme, but the resistance is 
still too high to be used in practical devices, although it can be used in Hall measurement 
at room temperature in place of In. Ta is probably a good candidate to form better contact 
than Ti. A diffusion barrier to prevent the formation of stable nitrides should also result in 
an improvement in device performance. Nb might be a good candidate for the diffusion 
barrier due to its slow diffusion rates in Ni. 

162 • 



163 

In the oxidized thin Ni/Au and Ni/Au/Ni schemes, more experimental work on 
sample cleaning, effects of annealing temperature and time, and O2 flowing rates is 
needed to find better processing parameters to decrease the contact resistance. On the 
other hand, while a low resistance of 4x10'^ Q-cm^ was reported, the work performed in 
this dissertation only showed specific contact resistance of 10 -10" Q-cm , which 
means the reported low value is not a typical result for this contact scheme. 

Also, this work showed that the contact pads have high resistances. Thick NiO 
layer on the contact would not allow much current to flow through the contacts, so those 
contacts with a thick NiO layer would be difficult to meet the specifications for 
applications requiring a thicker contact. More work is needed to decrease the resistivity 
of the NiO film. If the NiO can be made more conductive, the thickness of the cap-layer 
in the Ni/Au/Ni scheme can be increased further, while still keep the good electrical and 
optical properties. Doping with lithium is an option for this purpose. 

While this dissertation provides simple models for the high light transmittance 
and low specific contact resistance of the 50/50/50 scheme, they are only the beginning to 
address this complex problem. The effects of light absorption and scattering in the Au 
films need to be considered, as well as the effects of Ni cap-layer on reacfion kinetics. 
Computer simulation on the microstructure evolution in the annealed state should be 
helpful in obtaining a more in-depth and detailed understanding of the positive roles of Ni 
cap-layer in the thin Ni/Au contact scheme. Finally the role of the metallic Ni-rich layer 
in contact formation should be fiirther investigated. 









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BIOGRAPHICAL SKETCH 

Bo Liu was bom in Linshu, Shangdong province, in the People's Republic of 
China, on February 7, 1969. After completing 1 1 years of primary and secondary 
education in his home town, he moved to Harbin in the northeast region and enrolled in 
the Department of Metal Materials and Technology (now the College of Materials 
Science and Engineering) at the Harbin histitute of Technology (HIT) in 1987. In 1991, 
he was admitted to the graduate program at HIT with the highest grade. Under Professor 
Yang's guidance, he finished his research thesis on formability of intermetallic compound 
alloys in the spring of 1994 and worked continuously on the same research topic in 
Professor Yang's group. With a strong interest in the progress of semiconductor 
materials. Bo came to Gainesville to attend the University of Florida in 1997, majoring in 
electronic materials in the Department of Materials Science and Engineering. Under the 
supervision of Dr. Paul H. HoUoway, Bo worked on ohmic contact technology for 
gallium nitride and received his Ph.D in May 2001 . 



175 



I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 



^^^^^-^^^^-g^^ 



Paul H. Holloway, Chaimian 
Professor of Materials Science and 
Engineering 



I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 



Rolf E. Hummel 

Professor of Materials Science and 
Engineering 



I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philos^ 




Kevin S. Jones 

Professor of l^terials Science and 
Engineering 



I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 




Wolfgang M. Sigmund 
Assistant Professor of Materials Science 
and Engineering 



I certify that I have read this study and that in my opinion it conforms to 
acceptable standards of scholarly presentation and is fully adequate, in scope and quality, 
as a dissertation for the degree of Doctor of Philosophy. 




Fan Ren 

Professor of Chemical Engineering 



This dissertation was submitted to the Graduate Faculty of the College of 
Engineering and to the Graduate School and was accepted as partial fiilfillment of the 
requirements for the degree of Doctor of Philosophy. 



May 2001 




M. Jack Ohanian 

Dean, College of Engineering 



Winfred M. Phillips 
Dean, Graduate School 




./-^gs 



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